University of WollongongResearch Online
University of Wollongong Thesis Collection 2017+ University of Wollongong Thesis Collections
2018
Quality improvement in wire arc additivemanufacturingBintao WuUniversity of Wollongong
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Recommended CitationWu, Bintao, Quality improvement in wire arc additive manufacturing, Doctor of Philosophy thesis, School of Mechanical, Materials,Mechatronic and Biomedical Engineering, University of Wollongong, 2018. https://ro.uow.edu.au/theses1/473
QUALITY IMPROVEMENT IN WIRE ARC
ADDITIVE MANUFACTURING
A thesis submitted in fulfilment of
the requirements for the award of the degree
DOCTOR OF PHILOSOPHY
from
UNIVERSITY OF WOLLONGONG
By
BINTAO WU
B. Eng., M. Eng.
School of Mechanical, Materials, Mechatronics and Biomedical Engineering
Faculty of Engineering and Information Sciences
August, 2018
CERTIFICATION
I
CERTIFICATION
I, Bintao Wu, declare that this thesis, submitted in fulfilment of requirements for award
of Doctor of Philosophy at school of Mechanical, Materials, Mechatronics and Biomedical
Engineering, University of Wollongong, Australia, is wholly my own work unless otherwise
referenced and acknowledge. The document has not been submitted for qualifications at any
other or academic institution.
Type name: Bintao Wu
Signature:
Date: August, 2018
ACKNOWLEDGEMENTS
II
ACKNOWLEDGEMENTS
I wish to express my sincere appreciation and gratitude to my supervisor, Prof. Huijun Li, and
co-supervisor A/Prof. Zengxi Pan, and Dr. Dominic Cuiuri, for their valuable guidance,
support, inspiration and close supervision during my PhD candidature in University of
Wollongong (UOW). The assistance they provided in my studies over the last two years is the
key factor for the successful completion of this work.
I would like to express my warmest thanks to Dr. Donghong Ding, Chen Shen, and Yan Ma
for their valuable discussion, good friendship and encouragement.
I am deeply grateful to the staffs from the UOW Welding and Industrial Automation Research
Centre, and Australia institute for innovative materials center for their expertise and extensive
use of facilities. A special thanks to Matthew Franklin and Nathan Tarlinton for their support.
I would also like to convey sincere appreciation to all the members of the Faculty of
Engineering and Information Science at UOW and Engineering Enquiry Centre staff for
assisting in administration work.
I also appreciate the assistance from the China Scholarship Council (CSC), University of
Wollongong (UOW) for providing the scholarships to support my study.
Finally, I wish to express my special thanks to all my friends and to my family for their
support and encouragement during my PhD candidature career.
ABSTRACT
III
ABSTRACT
In recent years, significant progress has been made in the development of the Wire Arc
Additive Manufacturing (WAAM) processes due to its economically produced large-scale
metal components with relatively high deposition rate. As WAAM has evolved, a wide range
of materials have become associated with the processes and applications. Today, producing
high-quality WAAM products and improving their potential service life are still a challenge.
This dissertation focuses primarily on critical issues and methodologies to improve quality
performance of component fabricated by WAAM technique.
The research was the first attempt to study the WAAM process from complex thermal
behaviour point of view. It is found that due to the influences of thermal accumulation, the
interlayer’s surface oxidation, bead geometry, microstructural evolution, grain size, and
crystalline phase vary along the building direction of the as-fabricated wall, which creates
variations in arc shape, metal transfer behaviour, mechanical properties and fracture features.
Additionally, the corrosion behaviour within the WAAM-processed part is anisotropic due to
corresponding anisotropy in microstructure, phase structure, grain size and orientation. The
research provides a better understanding of the effects of heat accumulation behaviour on
deposition stability and material properties during WAAM process, which benefits future
process control, improvement, and optimization.
To achieve improved microstructure and mechanical properties, an innovative WAAM
process with forced interpass cooling using compressed CO2 was employed in this research.
ABSTRACT
IV
It is found that the forced interpass cooling is beneficial to additively manufactured
components, contributing to an appealing surface finish with less visible surface oxidation,
refined microstructure, improved hardness, enhanced strength and low distortion. Active
interpass cooling not only improves deposition properties, but also promotes geometrical
repeatability and also improved manufacturing efficiency through the reduction of dwell time
between layers.
A quality-based framework is proposed in the last section of the thesis, aiming to
produce high-quality and defect-free WAAM components. This thesis concludes that the wide
application of WAAM still presents many challenges, and these may need to be addressed in
specific ways for different materials in order to achieve an operational system in an
acceptable time frame. The integration of materials and manufacturing process to produce
defect-free and structurally-sound deposited parts remains a crucial effort into the future.
TABLE OF CONTENTS
V
TABLE OF CONTENTS
CERTIFICATION ................................................................................................................... I
ACKNOWLEDGEMENTS ................................................................................................... II
ABSTRACT ........................................................................................................................... III
TABLE OF CONTENTS ........................................................................................................ V
ABBREVITIATIONS AND SYMBOLS ............................................................................... X
LIST OF FIGURE.............................................................................................................. XIII
LIST OF TABLE ............................................................................................................... XVII
CHAPTER 1 INTRODUCTION ............................................................................................ 1
1.1 Background ................................................................................................................... 1
1.2 Objective of current research ........................................................................................ 2
1.3 Outline of thesis ............................................................................................................ 3
1.4 Original Research Contributions ................................................................................... 5
CHAPTER 2 LITERATURE REVIEW ................................................................................ 7
2.1 Wire Arc Additive Manufacturing (WAAM) systems .................................................. 7
2.2 Metals used in WAAM process ................................................................................... 10
2.2.1 Titanium alloys.................................................................................................. 10
2.2.2 Aluminum alloys and steel ................................................................................ 14
2.2.3 Ni-based superalloys ......................................................................................... 15
2.2.4 Other metals ...................................................................................................... 17
2.3 Common defects in WAAM-fabricated component.................................................... 18
2.3.1 Deformation and residual stress ........................................................................ 20
TABLE OF CONTENTS
VI
2.3.2 Porosity ............................................................................................................. 22
2.3.3 Crack and delamination .................................................................................... 23
2.4 Current methods for quality improvement in WAAM process ................................... 24
2.4.1 Post heat treatment ............................................................................................ 25
2.4.2 Interpass cold rolling......................................................................................... 26
2.4.3 Interpass heat sink ............................................................................................. 28
2.4.4 Peening and ultrasonic impact treatment .......................................................... 30
2.5 Summary and scope of this work ................................................................................ 31
CHAPTER 3 EXPERIMENTAL INSTRUMENTS AND METHODOLOGIES ............ 32
3.1 Materials ..................................................................................................................... 32
3.2 Experimental instruments ........................................................................................... 32
3.3 Process monitoring...................................................................................................... 33
3.4 Metallography ............................................................................................................. 35
3.4.1 Sample preparation for metallography .............................................................. 35
3.4.2 Stereo microscopy and optical microcopy (OM) .............................................. 37
3.4.3 Scanning electron microscopy (SEM) .............................................................. 38
3.4.4 X-ray diffraction (XRD) ................................................................................... 38
3.5 Mechanical properties ................................................................................................. 39
3.5.1 Tensile testing ................................................................................................... 39
3.5.2 Hardness testing ................................................................................................ 40
3.6 Corrosion resistance .................................................................................................... 41
3.7 Summary ..................................................................................................................... 42
CHAPTER 4 PROCESS STABILITY ................................................................................. 43
4.1 Introduction ................................................................................................................. 43
4.2 Experimental Procedures ............................................................................................ 44
4.2.1 Experimental setup............................................................................................ 44
4.2.2 Measurement of interpass temperature ............................................................. 45
4.3 Results and discussion ................................................................................................ 48
TABLE OF CONTENTS
VII
4.3.1 Interpass temperature and heat accumulation ................................................... 48
4.3.2 Bead appearance and geometrical features ....................................................... 50
4.3.3 Stability of weld pool and arc behaviour .......................................................... 53
4.4 Conclusion .................................................................................................................. 60
CHAPTER 5 MATERIAL PROPERTIES .......................................................................... 62
5.1 Introduction ................................................................................................................. 62
5.2 Experimental Procedures ............................................................................................ 63
5.2.1 Experiment setup .............................................................................................. 63
5.2.2 In-situ interpass temperature measurement ...................................................... 64
5.2.3 Heat accumulation calculation .......................................................................... 65
5.2.4 Material characterization techniques ................................................................ 66
5.3 Results and Discussion ............................................................................................... 67
5.3.1 Microstructure evolution ................................................................................... 67
5.3.2 Phases characteristics and transformation ........................................................ 72
5.3.3 Mechanical properties ....................................................................................... 73
5.3.4 Facture behaviours ............................................................................................ 75
5.4 Conclusion .................................................................................................................. 79
CHAPTER 6 CORROSION RESISTANCE ....................................................................... 81
6.1 Introduction ................................................................................................................. 81
6.2 Experimental Procedures ............................................................................................ 82
6.2.1 Sample and solution preparation ....................................................................... 82
6.2.2 Material analysis ............................................................................................... 84
6.2.3 Electrochemical measurements ......................................................................... 85
6.3 Results ......................................................................................................................... 86
6.3.1 Microstructural studies...................................................................................... 86
6.3.2. Electrochemical studies ................................................................................... 89
6.4 Discussion ................................................................................................................... 96
6.4.1. Microstructural evolution ................................................................................. 96
TABLE OF CONTENTS
VIII
6.4.2 Electrochemical evaluation ............................................................................... 98
6.5 Conclusion ................................................................................................................ 102
CHAPTER 7 THE FORCED INTERPASS COOLING .................................................. 103
7.1 Introduction ............................................................................................................... 103
7.2 Experimental Procedures .......................................................................................... 105
7.2.1 Experiment setup ............................................................................................ 105
7.2.2 Forced interpass cooling ................................................................................. 106
7.2.3 Material characterization techniques .............................................................. 107
7.3 Results ....................................................................................................................... 107
7.3.1 Deposition geometry ....................................................................................... 107
7.3.2 Surface oxidation ............................................................................................ 108
7.3.3 Macrostructure ................................................................................................ 110
7.3.4 Microstructure ................................................................................................. 113
7.3.5 Hardness .......................................................................................................... 114
7.3.6 Mechanical property ....................................................................................... 115
7.3.7 Fracture behaviour .......................................................................................... 117
7.4 Discussion ................................................................................................................. 118
7.5 Conclusion ................................................................................................................ 122
CHAPTER 8 DISTORTION CONTROL ......................................................................... 124
8.1 Introduction ............................................................................................................... 124
8.2 Experimental Procedures .......................................................................................... 126
8.2.1 Experiment setup ............................................................................................ 126
8.2.2 The deformation measurement........................................................................ 128
8.2.3 Modeling process ............................................................................................ 128
8.3 Results and discussion .............................................................................................. 132
8.3.1 Thermal behaviours during deposition ........................................................... 132
8.3.2 Geometrical features ....................................................................................... 134
8.3.3 Thermal distortion ........................................................................................... 136
TABLE OF CONTENTS
IX
8.4 Conclusion ................................................................................................................ 140
CHAPTER 9 CONCLUSION ............................................................................................. 142
9.1 A quality-based framework for WAAM process ....................................................... 142
9.2 Future perspective ..................................................................................................... 145
PUBLICATION .................................................................................................................... 147
REFERENCE ....................................................................................................................... 150
ABBREVIATIONS AND SYMBOLS
X
ABBREVIATIONS AND SYMBOLS
Abbreviations
WAAM Wire Arc Additive Manufacturing
GTAW gas tungsten arc welding
GMAW gas metal arc welding
CMT cold metal transfer
BPF Power-Bed Fusion
UOW University of Wollongong
CAD computer-aided design
3D 3-dimensional
ASTM American Society for Testing Materials
YS yield strength
UTS ultimate tensile strength
EI elongation
CMT-P cold metal transfer- pulse
CMT-ADV cold metal transfer- advanced
CO2 carbon dioxide
UIT ultrasonic impact treatments
AM Additive Manufacturing
LBM Laser Beam Melting
EBM Electron Beam Melting
IR infrared
H2O water
HNO3 concentrated nitric acid
ABBREVIATIONS AND SYMBOLS
XI
HF hydrofluoric acid
OM Optical Microcopy
SEM scanning electron microscopy
EDS energy dispersive X-ray spectroscopy
XRD X-ray diffraction
BM base material, base metal
HAZ heat affected zone
TIG tungsten inert gas
MIG metal inert gas
Symbols
beta phase, body-centred cubic structure in Ti6Al4V
alpha phase, face-centred cubic structure in Ti6Al4V
°C Celsius degree
g·cm-3 gram per cubic centimetre
at.% atom percent
𝜎𝑦 yield strength
𝜎0,𝑦 lattice frictional stress
d average grain diameter
L/min litre per minute
mm/min millimetre per minute
wt.% weight percent
rpm revolutions per minute
mm/s millimetre per seconds
° degree
ABBREVIATIONS AND SYMBOLS
XII
min-1 per minute
kV Kilovolts
mA milliampere
g gram
s-1 per seconds
LIST OF TABLE
XIII
LIST OF FIGURES
Figure 1-1 Different terminology of WAAM named by various research groups .................... 2
Figure 1-2 Outline of this dissertation ...................................................................................... 3
Figure 2-1 WAAM system design concepts, University of Wollongong .................................. 9
Figure 2-2 Optical microstructure of the deposited wall (from Lin, et al.[64]): (a) the bottom
region; (b) the middle region; (c) the top region. ............................................................. 12
Figure 2-3 Mechanical property comparison for Ti6Al4V parts fabricated using WAAM: (a)
yield and ultimate strength, (b) elongation. ...................................................................... 14
Figure 2-4 The major phases appearing in the as-deposited microstructure of Inconel 625
alloy (from Xu et al. [83]). ............................................................................................... 16
Figure 2-5 The correlation between materials and defects in WAAM process. ...................... 19
Figure 2-6 Schematic diagram of WAAM with cold rolling process [104] ............................ 27
Figure 2-7 Schematic diagram of the combined WAAM gas cooling process ....................... 29
Figure 2-8 Mechanical properties of Ti6Al4V part produced by WAAM with interpass
cooling using CO2 gas:(a) Hardness; (b) Tensile strength and elongation ....................... 30
Figure 3-1 The apparatus used in this work: (a) 200A-rated GTAW power source; (b) wire
feeder; (c) travel mechanism and trailing gas shield ........................................................ 33
Figure 3-2 High speed camera used in this study ................................................................... 34
Figure 3-3 IR pyrometer used in this study ............................................................................. 34
Figure 3-4 3D laser scanner used in this work ........................................................................ 35
Figure 3-5 Struers CitoPress-20 .............................................................................................. 36
Figure 3-6 Struers Tegrapol-21 ............................................................................................... 36
Figure 3-7 Leica M205A stereo microscopy ........................................................................... 37
Figure 3-8 Leica DMR optical microscopy ............................................................................ 37
Figure 3-9 JEOL JSM-7500 SEM ........................................................................................... 38
Figure 3-10 GBC MMA X-ray diffractometer ........................................................................ 39
Figure 3-11 MTS370 tensile machine ..................................................................................... 40
Figure 3-13 The Vickers innovative automatic testers ............................................................ 41
LIST OF TABLE
XIV
Figure 3-14 Electrochemical workstation ............................................................................... 42
Figure 4-1 Schematic illustration of GT-WAAM process ...................................................... 45
Figure 4-2 Schematic diagram of the temperature measurement system: (a) overall system (b)
measurement locations for the pyrometer and thermocouples (TC) ................................ 46
Figure 4-3 Pyrometer calibration curve for Ti6Al4V using fixed emissivity of 0.45 ............. 48
Figure 4-4 The variation in temperature and heat accumulation during fabrication............... 49
Figure 4-5 The bead appearance of (a) the first layer, (b) the top layer .................................. 51
Figure 4-6 Width of the build cross-section along the building height ................................... 52
Figure 4-7 The arc shape evolution in GT-WAAM of Ti6Al4V ............................................. 53
Figure 4-8 The schematic diagram of the changes in arc length and bead height for different
layers: (a) layer 1, (b) layer 5, (c) layer 10 and (d) layer 15 ............................................ 55
Figure 4-9 Metal transfer process at: (a) layer 1, (b) layer 5, (c) layer 15; (d) Interrupted
transfer occurring from layer 5 to 15 ................................................................................ 58
Figure 4-10 Metal transfer behaviour at different layers: (a) layer 1; (b) layer 5; (c) layer 15
.......................................................................................................................................... 58
Figure 4-11 Forces acting on droplet during Ti6Al4V GT-WAAM process .......................... 60
Figure 5-1 Schematic diagram of the GT-WAAM system ..................................................... 64
Figure 5-2The variation of temperature and heat accumulation in this study......................... 66
Figure 5-3 Cross-sectional macrograph of the deposited wall. ............................................... 68
Figure 5-4 Optical micrographs of the corresponding regions: a, b, c, d, e, f in Figure 5-3 .. 69
Figure 5-5Optical micrographs of selected samples in different horizontal planes: (a)
schematic of the sample locations; (b) base metal; (c) M1; (d) M2; (e) M3; (f) M4. ...... 71
Figure 5-6 Statistical distributions of the width of alpha lamellae for M1, M2 and M3 ........ 71
Figure 5-7 XRD spectrums of M1, M2, M3 and M4 .............................................................. 72
Figure 5-8 Mechanical test results of the selected samples in horizontal plane of the
deposited wall: (a) the locations of tensile samples; (b) load-displacement relationships;
(c) ultimate tensile strength and yield strength; (d) elongation and reduction of area ..... 74
Figure 5-9The fracture appearance of the tensile specimens .................................................. 77
Figure 5-10 High-magnification fractographs of corresponding tensile samples: (a) S1; (b)
S2; (c) S3; (d) S4; (e) S5; (f) S6; (g) S7; (h) S8. .............................................................. 78
LIST OF TABLE
XV
Figure 5-11 Axial crack in fracture surface of S7 ................................................................... 79
Figure 6-1 Schematic diagram of GT-WAAM system ........................................................... 83
Figure 6-2 Three-dimensional diagram of WAAM-fabricated Ti-6Al-4V wall showing
orientation of specimen planes ......................................................................................... 84
Figure 6-3 Optical micrographs of three selected specimens: (a) BM region; (b) HP region; (c)
VP region. ......................................................................................................................... 87
Figure 6-4 The XRD spectrums of three selected specimen from different planes ................ 88
Figure 6-5 Average hardness values of three groups of specimens ........................................ 89
Figure 6-6 Open circuit potential (vs. SCE) of as-received samples in 3.5% NaCl solution. 90
Figure 6-7 Potentiodynamic polarization plots of specimens ................................................. 92
Figure 6-8 Strong polarization curves for WAAM-fabricated Ti-6Al-4V and ASTM standard
Grade 5 alloy in 3.5wt% NaCl solution. The inset shows passive region from 0 V to 2 V.
.......................................................................................................................................... 93
Figure 6-9 EIS results of test samples: (a) Nyquist plots with inset showing the equivalent
circuit: (b) Bode plots. ...................................................................................................... 95
Figure 6-10 Schematic of microstructural evolution for WAAM-fabricated of Ti-6Al-4V: (a)
Continuous cooling diagram for Ti–6Al–4V β-solution treated at 1050°C for 30 min (Ms
temperature due to Majdic and Ziegler) [142], (b) HP regions, (c) VP regions. .............. 98
Figure 6-11 The comparison in corrosion resistance of WAAM-fabricated Ti-6Al-4V and
wrought base metal in 3.5% NaCl solution. ................................................................... 101
Figure 7-1 Schematic diagram of the GT-WAAM system ................................................... 105
Figure 7-2 Schematic diagram of the GT-WAAM deposition system with forced interpass
cooling ............................................................................................................................ 106
Figure 7-3 Schematic diagram of extracted tensile samples ................................................. 107
Figure 7-4 Effect of interpass temperature on wall dimensions ........................................... 108
Figure 7-5 The surface appearance of Ti6Al4V parts fabricated at different process
conditions: (a) 100°C; (b) 200°C; (c) 300°C; (d) Forced cooling with CO2 gas ........... 110
Figure 7-6 The cross-section morphology of Ti6Al4V component fabricated at : (a)100°C; (b)
200°C; (c) 300°C;(d) forced interpass cooling with CO2 gas ........................................ 112
Figure 7-7 The area ratio of parallel band regions and convex band regions for different
LIST OF TABLE
XVI
process conditions .......................................................................................................... 112
Figure 7-8 The microstructures of parts fabricated under different process conditions ....... 114
Figure 7-9 Hardness profiles of specimens: (a) Hardness distribution; (b) Average hardness
........................................................................................................................................ 115
Figure 7-10 The mechanical properties of obtained specimens ............................................ 117
Figure 7-11 High-magnification fractographs of tensile samples: (a) 100°C; (b) 200°C; (c)
300°C; (d) Interpass CO2 gas cooling ............................................................................ 118
Figure 8-1 Experimental setups of the GT-WAAM deposition system with forced interpass
cooling ............................................................................................................................ 126
Figure 8-2 Schematic of temperature measurement location for the thermocouple ............. 127
Figure 8-3 3D finite element mesh for one half of build plate ............................................. 129
Figure 8-4 The boundary conditions for the simulated model .............................................. 131
Figure 8-5 The temperature profiles of deposition (a) with natural cooling and (b) with active
interpass cooling ............................................................................................................. 133
Figure 8-6 The thermal history of one selected layer with different active interpass cooling
processes for specimens S2 to S6 ................................................................................... 134
Figure 8-7 Effect of forced interpass cooling on wall dimensions ....................................... 135
Figure 8-8 Simulation results for the distortion of produced samples: (a) S1; (b) S2; (c) S3;
(d) S4; (e) S5; (f) S6. ...................................................................................................... 137
Figure 8-9 The thermal distortion of as-deposited specimens: (a) longitudinal distortion and
(b) transverse distortion of half of substrate. .................................................................. 137
Figure 8-10 The distortion mechanism in WAAM-fabricated Ti6Al4V process .................. 138
Figure 9-1 A quality-based framework for WAAM process ................................................. 143
LIST OF TABLE
XVII
LIST OF TABLES
Table 2-1 Comparison of various WAAM techniques .............................................................. 8
Table 2-2 Metals typically used with WAAM process ............................................................ 10
Table 2-3 Mechanical properties of titanium alloys fabricated from various WAAM processes
.......................................................................................................................................... 13
Table 2-4 Tensile properties of WAAM-fabricated aluminium alloy (2219) .......................... 15
Table 2-5 Mechanical properties of various Ni-based superalloys using different WAAM
processes ........................................................................................................................... 17
Table 2-6 Mechanical properties of other metallic materials fabricated using WAAM process
.......................................................................................................................................... 18
Table 2-7 Tendency of various defects in WAAM fabricated parts ........................................ 20
Table 2-8 Material properties of component fabricated using WAAM with interpass cold
rolling ............................................................................................................................... 27
Table 2-9 The distribution of porosity in aluminum part using WAAM with cold rolling. .... 28
Table 3-1 Chemical composition of Ti-6Al-4V (wt.%)........................................................... 32
Table 3-2 Sample preparation procedures for microstructural analysis. ................................. 36
Table 3-3 Dimensions for tensile specimen. ........................................................................... 40
Table 4-1 Process parameters for GT-WAAW ........................................................................ 45
Table 4-2 Detailed process data for heat accumulation calculation. ....................................... 50
Table 5-1 Process parameters for WAAM deposition ............................................................. 64
Table 6-1 Process parameters for GT-WAAM deposition ...................................................... 83
Table 6-2 Electrochemical kinetics parameters for test samples ............................................. 92
Table 6-3 EIS parameters of equivalent circuit for test samples ............................................. 96
Table 7-1 Process parameters for WAAM deposition ........................................................... 106
Table 8-1 Parameter design for cooling process ................................................................... 128
Table 8-2 Welding and deposition process parameters ......................................................... 128
CHAPTER 1 INTRODUCTION
1
Chapter 1 INTRODUCTION
1.1 Background
In recent years, wire arc additive manufacturing (WAAM) has increasingly attracted
attention from the industrial manufacturing sector due to its ability to create large metal
components with high deposition rate, low equipment cost, high material utilization, and
consequent environmental friendliness. The origin of the WAAM process can be traced back
to 1925s when Baker [1] proposed to use an electric arc as the heat source with filler wires as
feedstock materials to deposit metal ornaments. Since then, consistent progress has been
made on the development of this technology, particularly in the last 10 years; the WAAM
technique bears various nomenclatures by different research institutions worldwide [2-24], as
shown in Figure 1-1. Today, WAAM has become a promising fabrication process for various
engineering materials such as titanium, aluminium, nickel alloy and steel. Compared to
traditional subtractive manufacturing, the WAAM system can reduce fabrication time by
40-60% and post-machining time by 15-20% depending on the component size[25]. For
instance, recent breakthrough in WAAM technology has made it possible to fabricate landing
gear rib with a saving of approximately 78% in raw material when compared with the
traditional subtractive machining process[26].
CHAPTER 1 INTRODUCTION
2
Figure 1-1 Different terminology of WAAM named by various research groups
Due to the highly complex nature of WAAM, many different aspects of the process need
to be studied, including processes development, material quality and performance, path
design and programming, process modelling, process monitoring and online control[27].
Several WAAM review papers have been published by leaders in the field, covering the
state-of-the-art systems, design, usage, in-situ process monitoring, in-situ metrology and
in-process control and sensing [26, 28-32].
1.2 Objective of current research
The main objective of this research is to investigate the effects of thermal accumulation
on the deposition stability and material properties during WAAM process with the aim of
providing strategies for further process control and optimization. Furthermore, to avoid the
generation of the defects during deposition and to improve manufacturing efficiency, an
innovative WAAM process with active interpass cooling using compressed CO2 gas is
proposed. Based on the research work, a quality-based framework will be presented in
CHAPTER 1 INTRODUCTION
3
conclusion, aiming to achieve high-quality and defect-free WAAM components. The research
tasks completed in this dissertation are now listed in Figure 1-2.
Figure 1-2 Outline of this dissertation
1.3 Outline of thesis
This thesis is organized as follows:
Chapter 2 reviews the microstructure and mechanical properties of various metals,
including titanium and its alloys, aluminum and its alloys, Ni-based alloy, steel and other
intermetallic materials fabricated by the various WAAM processes. The common defects that
have been found to occur for different materials are also summarized. The current methods
for both in-process and post-process quality improvement and defect reduction are
introduced.
Chapter 3 introduces the experiment instruments and methodologies that have been used
in this study.
Chapter 4 investigates the influence of heat accumulation on bead formation, arc
CHAPTER 1 INTRODUCTION
4
stability, metal transfer behaviour during the manufacture of Ti6Al4V with the gas tungsten
wire arc additive manufacturing (GT-WAAM) using localized gas shielding. An infrared
pyrometer is used to measure the in-situ interpass temperature and arc stability and metal
transfer behaviour are monitored by means of a high-speed camera.
Chapter 5 further investigates the effects of heat accumulation on microstructure and
mechanical properties of additively manufactured Ti6Al4V parts by means of optical
microscopy (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), energy
dispersive spectrometer (EDS) and standard tensile tests, aiming to explore the quality
performance of fabricating Ti6Al4V parts by GT-WAAM using localized gas shielding.
Chapter 6 investigates the relationships between corrosion resistances, microstructure
and phase composition of Ti-6Al-4V components that have been fabricated with the gas
tungsten wire arc additive manufacturing (GT-WAAM) process, through the use of
electrochemical corrosion testing, optical microscopy, X-ray diffraction and hardness testing,
trying to provide a strategy for the anisotropic corrosion behaviour of additive manufactured
part.
Chapter 7 proposes an innovative manufacturing process with forced interpass cooling
using compressed CO2 to achieve improved microstructure and mechanical properties. The
effects of various interpass temperatures and rapid forced cooling on deposition geometry,
thermal distortion, surface oxidation, microstructural evolution, and mechanical properties of
the fabricated part were investigated.
Chapter 8 conducts a comparative analysis of simulation and experimental investigation
on the effects of forced interpass cooling on the thermal state during deposition process, and
CHAPTER 1 INTRODUCTION
5
further to characterize the geometry and thermal distortion of build part.
Chapter 9 gives a discussion on improving quality of WAAM fabricated parts through
process selection, feedstock optimization, process monitoring and control and post-process,
including proposals for future research directions.
1.4 Original Research Contributions
The main contributions of the dissertation are threefold.
(1) In-depth study of the heat accumulation effects on the process stability and
material properties of WAAM-fabricated component.
Complex thermal behaviour during fabrication plays an import role in the process
stability, geometrical formation and mechanical properties of components manufactured
using WAAM technology. Owing to the influences of heat accumulation, the interlayer
surface oxidation, bead geometries, microstructural evolution, grain size, and crystalline
phase vary along the building direction of the as-fabricated wall, which leads to the variations
in arc shape, metal transfer behaviour, mechanical properties and fracture features. A better
understanding of the relations between thermal accumulation and process stability, material
properties will benefit future process control, improvement, and optimization.
(2) A newly developed active interpass cooling technique for WAAM to achieve
improved component quality.
An innovative wire arc additive manufacturing (WAAM) process with forced interpass
cooling using compressed CO2 was employed in this dissertation. Forced interpass cooling
using compressed CO2 gas is easily implemented, and is beneficial to additively
manufactured Ti6Al4V components, contributing to an appealing surface finish with less
CHAPTER 1 INTRODUCTION
6
visible surface oxidation, refined microstructure, improved hardness, enhanced strength and
low distortion. Furthermore, it significantly promotes manufacturing efficiency through a
sharp reduction of dwell time between the deposited layers.
(3) The development of a quality-based framework for WAAM components
Based on an in-depth understanding of various materials, ideal process setup, in-process
parameter control and post processing, a quality-based framework is proposed, for producing
high-quality and defect-free components. Three main aspects are primarily considered:
feedstock optimization, manufacturing process, and post-process treatment.
CHAPTER 2 LITERATURE REVIEW
7
Chapter 2 LITERATURE REVIEW
This chapter reviews the microstructure and mechanical properties of various metals,
including titanium and its alloys, aluminum and its alloys, Ni-based alloy, steel and other
intermetallic materials fabricated by the various WAAM processes. The common defects that
have been found to occur for different materials are also summarized. The current methods
for both in-process and post-process quality improvement and defect reduction are introduced.
Finally, a discussion is given on improving quality of WAAM fabricated parts through
process selection, feedstock optimization, process monitoring and control and post-process,
including proposals for future research directions.
2.1 Wire Arc Additive Manufacturing (WAAM) systems
Depending on the nature of the heat source, there are commonly three types of WAAM
processes: Gas Metal Arc Welding (GMAW)-based[33], Gas Tungsten Arc Welding
(GTAW)-based[2] and Plasma Arc Welding (PAW)-based[3]. As listed in Table 2-1, specific
class of WAAM techniques exhibit specific features. The deposition rate of GMAW-based
WAAM is 2-3 times higher than that of GTAW-based or PAW-based method. However, the
GMAW-based WAAM is less stable and generates more weld fume and spatter due to the
electric current acting directly on the feedstock. The choice of WAAM technique directly
influences the processing conditions and production rate for a target component.
CHAPTER 2 LITERATURE REVIEW
8
Table 2-1 Comparison of various WAAM techniques
WAAM Energy source Features
GTAW-based GTAW
Non-consumable electrode; Separate wire feed process;
Typical deposition rate: 1-2kg/hour;
Wire and torch rotation are needed;
GMAW-based
GMAW
Consumable wire electrode;
Typical deposition rate 3-4kg/hour;
Poor arc stability, spatter;
Cold metal transfer
(CMT)
Reciprocating consumable wire electrode;
Typical deposition rate: 2-3kg/hour;
Low heat input process with zero spatter, high process tolerance;
Tandem GMAW
Two consumable wires electrodes;
Typical deposition: 6-8kg/hour;
Easy mixing to control composition for intermetallic materials
manufacturing ;
PAW-based Plasma
Non-consumable electrode; Separate wire feed process;
Typical deposition rate 2-4kg/hour;
Wire and torch rotation are needed;
Most WAAM systems use articulated industrial robot as the motion mechanism. Two
different system designs are available. The first design uses an enclosed chamber to provide a
good inert gas shielding environment similar to laser Power-Bed Fusion (PBF) system. The
second design uses existing or specially designed local gas shielding mechanism with the
robot positioned on a linear rail to increase the overall working envlop. It is capable of
fabricating very large metal structure up to a meter in dimension. Figure 2-1 shows an
example of this design of WAAM system, used for the research and development at the
University Wollongong (UOW).
CHAPTER 2 LITERATURE REVIEW
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Figure 2-1 WAAM system design concepts, University of Wollongong
Fabricating a part using WAAM involves three main steps: process planning, deposition,
and post processing. For a given CAD model, 3D slicing and programming software
generates the desired robot motions and welding parameters for the deposition process, aimed
at producing defect-free fabrication with high geometrical accuracy [22, 23, 34]. Based on the
welding deposition model for the specific material being used to fabricate the component, the
3D slicing and programming software offer automated path planning and process
optimization to avoid potential process-induced defects [35-38]. During fabrication, the robot
and external axis provide accurate motion for the welding torch to build up the component in
a layer-by-layer fashion. Advanced WAAM systems can be equipped with various sensors to
measure welding signals [39], deposited bead geometry [40], metal transfer behaviours[41]
and interpass temperature[33, 42, 43], thereby supporting in-process monitoring and control
to achieve higher product quality. This is an area of current and future research interest, with
CHAPTER 2 LITERATURE REVIEW
10
the potential for significantly improving WAAM process performance.
2.2 Metals used in WAAM process
WAAM processes use commercially available wires which are produced for the welding
industry and available in spooled form and in a wide range of alloys as feedstock materials.
Table 2-2 indicates the commonly used alloys and their various applications in WAAM.
Manufacture of a structurally sound, defect free, reliable part requires an understanding of the
available process options, their underlying physical processes, feedstock materials, process
control methods and an appreciation of the causes of the various common defects and their
remedies. This section reviews the metals that are commonly used in WAAM, with a
particular emphasis on the microstructure and mechanical properties of the additively
manufactured alloys.
Table 2-2 Metals typically used with WAAM process
Applications
Alloys
Ti-based Al-based Steel-based Ni-based Bimetallic
Aerospace [26] [44] - [45] [46]
Automotive - [47] [48] - [49]
Marine [50, 51] - [52] - -
Corrosion resistance [53] - - [54] [55]
High temperature [56] - - [57] [58]
Tools and molds - - [59, 60] - -
2.2.1 Titanium alloys
Titanium alloys have been widely studied for application of additive manufacturing in
aerospace components due to their high strength-to-weight ratio and inherently high material
CHAPTER 2 LITERATURE REVIEW
11
cost. There are increasing demands for more efficient and lower cost alternative to the
conventional subtractive manufacturing methods, which suffers very low fly-to-buy ratios for
many component designs. There exists many business opportunities for the WAAM process,
particularly for large-sized titanium components with complex structures [26].
It is well accepted that the microstructure of the product depends on its thermal history
during the fabrication process. The distinctive thermal cycle, which involves an alternating
heating and cooling [61, 62], produces meta-stable microstructures and inhomogeneous
compositions in the fabricated part[63]. For example, Baufeld et al [62] investigated the
microstructures of Ti6Al4V fabricated using a GTAW-based WAAM system, and found two
distinctive regions on the as-build wall. In the bottom region, where alternating bands are
perpendicular to the build direction, a basket Widmanstätten structure with α phase lamellae
is present, while in the top region, where no such bands appear, needle-like precipitate is the
main structure. Similar microstructural evolution has also been observed in PAW-based
process. Lin, et al [64, 65] reported a graded microstructure along the build direction and
identified the martensite α’ structure, Widmanstätten structure and basket-wave structure
from the bottom to the top region of the fabricated component, as shown in Figure 2-2. An
epitaxial growth of β grains with discrete direction is also observed along the build direction
owing to thermal gradient [66], commonly seen in additively manufactured titanium alloy
components.
CHAPTER 2 LITERATURE REVIEW
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Figure 2-2 Optical microstructure of the deposited wall (from Lin, et al.[64]): (a) the bottom
region; (b) the middle region; (c) the top region.
Table 2-3 summarizes the microstructure and mechanical property data (tensile strength,
yield strength and elongation) of Ti6Al4V samples fabricated using various WAAM
technologies [62, 64, 67-73]. The as-forged and as-cast minimum specifications from ASTM
standards are also listed for comparison. As shown in Figure 2-3, the tensile property of
asfabricated Ti6Al4V samples is close to that of wrought Ti6Al4V and exceeds that of cast
Ti6Al4V as specified by ASTM standards. In addition, WAAM fabricated Ti6Al4V samples
show anisotropic properties with lower strength and higher elongation values in the build
direction (Z) compared to deposition direction (X), which is mainly attributed to the grain
size of α lamellae and the orientation of the elongated prior β grains.
CHAPTER 2 LITERATURE REVIEW
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Table 2-3 Mechanical properties of titanium alloys fabricated from various WAAM processes
Process Condition Microstructure YS[MPa] UTS[MPa] EL[%] Reported by
Cast / / 758 860 >8 ASTM F1108
Wrought / / 860 930 >10 ASTM F1472
GTAW
AF Columnar prior β grains +
Widmanstätten α/β /
929 ± 41a
965 ± 39b
9±1.2a
9±1b Baufeld et al. [13]
AF α phase lamella
basket weave structures /
939 ± 24 a
1033 ±32 b
16 ± 3a
7.8 ± 2.3b
Baufeld et al.[67] HT (600°C/4h/FC) lamellar structure / 972 ± 41 a
977± 14 b
12.5 ± 2.5a
6 ± 3b
HT (834°C/2h/FC) lamellar structure / 931 ±19 a
971 ± 28 b
21 ± 2a
14 ± 2b
AF Widmanstätten α + banded
coarsened lamella α
803 ± 15 a
950 ± 21 b
918 ±17 a
1033 ±19 b
14.8 a
11.7 b Wang et al.[74]
AF / 861 ± 14 a
892 ± 31 b
937 ± 21 a
963 ± 22 b
16.5 ± 2.7a
7.8 ± 2b
Brandl et al.[71] HT (600°C/2h/FC) / 891 ± 16 a
915 ± 14 b
976 ± 35 a
981 ± 8 b
11.6 ± 2.4a
6.6 ± 2.6b
AN(834°C/2h/FC) / 856 ± 21 a
893 ± 24 b
931 ± 17 a
962 ± 29 b
20.4 ± 1.8a
13.5 ± 2b
Plasma AF
600°C / 840°C
Widmanstätten α/β +
Columnar β grains / / / Martina et al.[75]
Pulsed-PAM AF Prior columnar β +
Martensite α’ 909 ±13.6b 988±19.2b 7±0.5b Lin et al.[65]
PAM AF
Prior columnar β +
martensite α’+ fine
basket-weave structure
877 ± 18.5 b 968 ± 12.6 b 11.5 ± 0.5 b Lin et al. [64]
AF: as fabricated, HT: heat treated, AN: annealed; a: In build direction, b: Orthogonal to build direction;
(a)
CHAPTER 2 LITERATURE REVIEW
14
(b)
Figure 2-3 Mechanical property comparison for Ti6Al4V parts fabricated using WAAM: (a)
yield and ultimate strength, (b) elongation.
2.2.2 Aluminum alloys and steel
Although fabrication trials for many different series of aluminum alloys, including Al-Cu
(2xxx)[76], Al-Si (4xxx)[77] and Al-Mg (5xxx)[41] have been successfully carried out, the
commercial values of WAAM is mainly justifiable for large and complex thin-walled
structures, since cost of manufacturing small and simple aluminum alloy components using
conventional machining processes is low[78]. Using WAAM to fabricate steel is unpopular
for the same reason although it is the most commonly used engineering material[79]. Another
reason for the poor commercial application of WAAM in aluminium is that some series of
aluminum alloys, such as Al 7xxx and 6xxx, are challenging to weld[80] due to turbulent
melt pool and weld defects, which frequently occur during the deposition process.
In general, as-deposited additively manufactured aluminum alloy parts have inferior
mechanical properties compared to those machined from billet material. In order to achieve
CHAPTER 2 LITERATURE REVIEW
15
higher tensile strength, most of the as-deposited aluminum parts undergo post-process heat
treatment to refine the microstructure. Table 2-4 lists the yield strength (YS), ultimate tensile
strength (UTS), and elongation of WAAM-fabricated 2219 aluminium alloy samples. Due to
the uniform distribution of large diamond particles within the microstructure, the sample
exhibits lower UTS and YS than that of the wrought part specified by ASTM standard.
However, after heat treatment, significant improvement beyond ASTM standard can be
observed in both strength and elongation as a consequence of the grain refinement [44].
Table 2-4 Tensile properties of WAAM-fabricated aluminium alloy (2219)
Materials Process Condition Microstructure YS[MPa] UTS[MPa] EL[%] Reported by
Al6.3Cu
Wrought(2219) T851 / 267 390 > 4 ASTMB211M[81]
CMT
AF Fine dendrites +
equiaxed grains
128 ± 2a
133 ± 5b
262 ± 4a
264 ± 2b
15.8 ± 0.3a
18.6 ± 1.5b Gu et al. [44]
HT(T6) Homogeneous
dispersed θ precipitates
305 ± 6a
333 ± 6b
458 ± 3a
466 ± 3b
13.6 ± 0.9a
14b
AF: as fabricated, HT: heat treated; a: In build direction, b: Orthogonal to build direction;
2.2.3 Ni-based superalloys
Ni-based superalloys are the second most popular material studied by the additive
manufacturing community research community after titanium alloys, mainly due to their high
strengths at elevated temperatures and high fabrication cost using traditional methods.
Nickel-based superalloys are widely applied in aerospace, aeronautical, petrochemical,
chemical and marine industries due to their outstanding strength and oxidation resistance at
temperature above 550 . To date, various Nickel-based superalloys, including Inconel 718
and Inconel 625 alloy have been studied after WAAM processing.
The microstructure of WAAM fabricated Inconel 718 parts generally consists of large
CHAPTER 2 LITERATURE REVIEW
16
columnar grains with interdendritic boundaries delineated by small Laves phase precipitates
and MC carbides[82]. Xu et al. [83] reported that columnar dendrite structures decorated with
a large amount of Laves phase, MC carbides and Ni3Nb are also present in WAAM-fabricated
Inconel 625 parts, as shown in Figure 2-4. It is worth noting that the microstructure can be
refined to smaller dendritic arm spacing, less niobium segregation and discontinuous Laves
phase in the interdendritic regions using post-process heat treatments, which are beneficial to
the mechanical properties.
Figure 2-4 The major phases appearing in the as-deposited microstructure of Inconel 625
alloy (from Xu et al. [83]).
Table 2-5 lists the mechanical properties of several Nickel-based superalloys fabricated
using the WAAM process. For GMAW-based WAAM-fabricated Inconel 718 alloy, the yield
and ultimate tensile strength is 473 ± 6 MPa and 828 ± 8 MPa respectively. These values lie
between the minimum values specified by ASTM for wrought and cast materials, whereas the
elongation is much lower than the standards for both wrought and cast conditions. As for
WAAM-fabricated Inconel 625 alloy, the YS, UTS and elongation all meet the requirement
CHAPTER 2 LITERATURE REVIEW
17
set by ASTM for cast materials and are slightly lower than those for wrought material.
Table 2-5 Mechanical properties of various Ni-based superalloys using different WAAM
processes
Materials Process Condition Microstructure YS[MPa] UTS[MPa] EL[%] Reported by
IN 718 GMAW AF Nb precipitates +
dendritic structure 473 ± 6 828±8 28 ± 2 Baufeld. [82]
IN 625
Cast / / 350 710 48 ASM[84]
Wrought / / 490 855 50 ASM[84]
PPAM AF Laves phase + columnar
dendrite structure
438 ± 38a
423 ± 22b
721 ± 32a
718 ± 19b
48.6 a
49.2 b Xu et al. [83]
PPAW
PPAW
GTAW
AF Laves phase + MC
carbides + δ-Ni3Nb 449 726 43
Xu et al. [85]
Xu et al. [85]
Wang et al. [86]
IC Laves phase +
NbC carbides 480 771 50
HT(980°C /STA) Coarser Laves particles +
Nb precipitates 469 802 42
CMT AF Nb,Mo precipitates +
dendrite structure /
684± 23 a
722 ± 17 b
40.13 ±3.7 a
42.27±2.4 b Ding et al. [87]
NiAlCu CMT AF Widmanstätten α +
martensite 350 ± 17 667 ± 15 29 ± 2.6 Ding et al. [87]
AF: as fabricated, HT: heat treated, IC: interpass cooling, STA: direct aging;
a: In build direction, b: Orthogonal to build direction,
2.2.4 Other metals
Other metals have also been investigated for potential fabrication using WAAM, such as
magnesium alloy AZ31 for automotive applications[88], Fe/Al intermetallic compound [89,
90] and Al/Ti [91, 92] compound, as well as bimetallic steel/nickel [93] and steel/bronze [94]
parts for aeronautic industry,. The detailed mechanical properties of these materials fabricated
using WAAM are listed in Table 2-6. Most of this research has focused on determining the
microstructural and mechanical properties of samples taken from simple straight-walled
structures, rather than developing a process to fabricate functional parts. Manufacturing
CHAPTER 2 LITERATURE REVIEW
18
intermetallic parts with accurate pre-designed composition still poses major challenges for
the WAAM process.
Table 2-6 Mechanical properties of other metallic materials fabricated using WAAM process
Materials Process Condition Microstructure YS[MPa] UTS[MPa] EL[%] Reported by
Ti/Al GTAW
AF Interdendritic γ structures +
fully lamellar colonies
424 ± 30a
474 ± 17b
488 ± 50a
549 ± 23b
0.35a
0.5b
Ma et al. [24] HT(1200°C/24h) Full γ microstructure 425 ± 15a
471 ± 14b
412 ± 11a
472 ± 11b
1.1a
1.1b
HT(1060°C/24h) Fully lamellar structure 487 ± 14a
486 ± 12a
569 ± 12a
590 ± 4b
0.30a
0.45b
Cu/Al GTAW AF copper twinning 63 ± 2.1a 231 ± 2.5a 63 ± 4.0a Dong et al. [95]
Fe/Al GTAW AF Columnar Fe3Al grains 847 ± 2b 944 ± 19 b 3.2 ± 0.1 b Shen et al. [89]
Fe/Ni GMAW AF austenite , δ-ferrite in low part;
dendritic in up part / 565±15 a /
Takeyuki
et al. [93]
Fe/Cu GMAW AF the mixture of α-Fe and ε-Cu / 305a 48.5 Liu et al. [94]
AF: as fabricated, HT: heat treated; a: In build direction, b: Orthogonal to build direction.
2.3 Common defects in WAAM-fabricated component
Although the mechanical properties of components fabricated by WAAM are in many
cases comparable to those of their conventionally processed counterparts, there are however
some AM processing defects that must be addressed for critical applications. Porosity, high
residual stress levels, and cracking, must be avoided, particularly for parts exposed to
extreme environments where these defects lead to failure modes such as high temperature
fatigue. Defects in WAAM can occur for various reasons, such as poor programming strategy,
unstable weld pool dynamics due to poor parameter setup, thermal deformation associated
with heat accumulation, environmental influence (such as gas contamination) and other
machine malfunctions. As shown in Figure 2-5, certain materials tend to be vulnerable to
CHAPTER 2 LITERATURE REVIEW
19
specific defects. For example, severe oxidization for titanium alloys, porosity for aluminum
alloys, poor surface roughness in steel as well as severe deformation and cracks in bimetal
components have been found to typically occur. Table 2-7 lists the major defects that are
commonly present in components fabricated using current WAAM techniques. The details of
these common defects and their relationship to the materials will be discussed this section.
Figure 2-5 The correlation between materials and defects in WAAM process.
CHAPTER 2 LITERATURE REVIEW
20
Table 2-7 Tendency of various defects in WAAM fabricated parts
Material Process Defect of feature
Ref. Porosity Cracking Delamination Oxidation Substrate adherence Surface finish
Ti6Al4V
TIG No No No Light Good Smooth [13]
Plasma No No No No Good Smooth [75]
CMT No No No Light Good Smooth [96]
DCEP-GMAW No No No Light Medium Poor [96]
H08Mn2Si steel DE-GMAW low No No No Good Waviness [18]
Copper-coated steel GMAW No No No Light Good Medium-rough [97]
ER4043 Al alloy CMT High No No Light Good Smooth [77]
VP-GTAW No No No No Good Medium-rough [98]
AA2319 Al alloy CMT High No No No Good Smooth [99]
CMT-PADV No No No No Good Smooth [99]
5356 Al alloy VP-GTAW No Yes No No Good Smooth [100]
Inconel 625 PPAD High Yes No No Good Smooth [83]
GTAW No No No No Good Smooth [101]
Inconel 718 GMAW Medium Yes Yes No Good Smooth [14]
AZ31 Mg alloy PMIG No No No Light Medium Medium-rough [88]
Nickel-Al-Cu CMT No No No No Good Smooth [87]
Steel–bronze bimetal GMAW No No No No Good Smooth [94]
Steel- nickel bimetal GMAW No No No No Good Medium-rough [93]
Intermetallic Fe/Al GTAW High Yes No Serious Medium Medium-poor [90]
Intermetallic Al/Ti GTAW Low Yes No No Good Rough [24]
Intermetallic Al/Cu GTAW No No Yes Light Poor Rough [102]
2.3.1 Deformation and residual stress
Like other additive manufacturing process, distortion and residual stress are inherent to
the WAAM process and it is impossible to completely avoid its generation. The residual
stress can lead to distortion of the part, loss of geometric tolerance, delamination of layers
during deposition, as well as deterioration of fatigue performance and fracture resistance of
the additively manufactured components. Therefore, control and minimization of deformation
and residual stress is a key area if research.
CHAPTER 2 LITERATURE REVIEW
21
Various types of deformation appear in WAAM fabricated parts, including longitudinal
and transvers shrinkage, bending distortion, angular distortion and rotational distortion [103].
The distortions are caused by thermal expansion and shrinkage of the part during repeated
melting and cooling processes, which is particularly an issue for large thin walled structures
[104]. Residual stress is the stress that remains in the material when all external loading
forces are removed. If the residual stress is sufficiently high, it can be a critical influential
factor in the mechanical properties and fatigue performance of the fabricated component. If
the residual stress exceeds the local UTS of the material, cracking will take place, while if it
is higher than the local YS but lower than UTS, warping or plastic deformation will occur
[105]. Ding et al. [33] found that the residual stress uniformly distributes across the WAAM
deposited wall, and the residual stress in the preceding layer has little effect on the future
layers. After release of clamping, however, the internal stress is redistributed with a much
lower value at the top of integral part than at the interface to the substrate, resulting in
bending distortion of the component. Path planning also involves the distortion and residual
stress evolution in WAAM process. If appropriate deposition path designs, it will help in the
significant improvement in these defects, especially in large metal fabrication. A detailed
overview of the residual stress origin would exceed the scope of this article.
Among all WAAM engineering materials, bimetal components exhibit high levels of
residual stress and deformation due to the material thermal expansion difference. Hence,
accurate interpass temperature control is needed when bimetal materials are used.
WAAM-fabricated Inconel alloy has comparatively lower residual stresses levels, but it is
more susceptible to process defects such as delamination, buckling and warping, since its
CHAPTER 2 LITERATURE REVIEW
22
residual stress is usually higher than the yield stress [106]. Other comparatively softer
materials, such as aluminum alloys, easily respond to deformation defects due to their high
thermal expansion coefficients. A better understanding pf the effect of material characteristics
in WAAM processing is needed for controlling residual stress and deformation during
deposition.
Deformation and residual stress are associated with many process parameters, such as
welding current, welding voltage, feeding speed, ambient temperature, shielding gas flow rate,
etc. There is still a lack of systematic methods for controlling defects through optimized
selection or manipulation of these parameters. Fortunately, several post-process treatments
that have been proven to effectively mitigate residual stress and deformation, and these will
be discussed in later section.
2.3.2 Porosity
Porosity is another common defect in WAAM processing that needs to be minimized
due to adverse effects on mechanical properties [107]. Firstly, porosity will lead to a
component with low mechanical strength by damage from micro-cracks, and secondly, it
usually brings low fatigue property to deposition via spatially with different size and shape
distribution. In general, this type of defects are mainly classified as either raw
material-induced [108] or process-induced[109].
The WAAM raw material, including as-received wire and substrate, often has a degree
of surface contamination, such as moisture, grease and other hydrocarbon compounds that
may be difficult to completely remove. These contaminants can be easily absorbed into the
CHAPTER 2 LITERATURE REVIEW
23
molten pool and subsequently generate porosity after solidification. Among common
engineering materials, aluminum alloy is the most susceptible to this defect as the solubility
of hydrogen in solid and liquid (0.036 against 0.69 cm3/100 g at a melting point of 660,
respectively) is significantly different [110]. Even small amounts of dissolved hydrogen in
the liquid state may exceed the limit of solubility after solidification, resulting in porosity
[111]. Therefore, the cleanliness of raw materials is critical, especially for aluminum alloys.
Process-induced porosity is usually non-spherical, and mainly caused by poor path
planning or an unstable deposition process. When the deposition path is complex, or the
manufacturing process is changeable, insufficient fusion or spatter ejection is easily produced,
creating gaps or voids in these influenced regions.
To control porosity, the following methods can be adopted: (1) an AC GMAW-based
process or CMT-PADV based process (cold metal transfer pulsed advanced, a controlled
short-circuiting GMAW transfer process) is preferred, especially for aluminum; (2) the
highest quality shielding gas, tight gas seals, non-organic piping and short pipe lengths are
highly recommended; (3) the wire and substrate surfaces are as clean as possible before
fabrication; (4) high quality feedstock should be used; (5) the deposited bead shape needs to
be optimized; (6) the thermal profile during processing should be monitored and controlled;
(7) post-deposition treatment, such as interpass rolling can be applied.
2.3.3 Crack and delamination
Similar to residual stress and deformation, cracking and delamination not only involves
the thermal signature of the manufacturing process, but also relates to the material
characteristics of the deposit. Ordinarily, the crack is categorised as either a solidification
CHAPTER 2 LITERATURE REVIEW
24
crack or grain boundary crack within the WAAM component [105]. The former type of crack
depends mainly on the solidification nature of the material and is usually caused by the
obstruction of solidified grain flow or high strain in the melt pool [112]. Grain boundary
cracking often generates along the grain boundaries due to the differences between boundary
morphology and potential precipitate formation or dissolution [113]. Delamination or
separation of adjacent layers takes place due to incomplete melting or insufficient re-melting
of the underlying solid between layers. Generally, this deficiency is visible and cannot be
repaired by post-process treatment. In order to prevent this defect, preprocess treatment such
as preheating of the substrate needs to be considered.
Bimetal material combinations, such as Al/Cu, Al/Ti and Al/Fe, are quite susceptible to
cracking and delamination when fabricated with the WAAM process. The dissimilar metals
have large differences in their mutual solubility and chemical reactivity so that the
intermetallic phase-equilibrium is freely broken, thus inducing crack growth along grain
boundaries. Also, Inconel alloy readily generates solidification cracking issues due to the
existence of liquid film at terminal solidification [101]. Both of these material types should
receive particular attention to avoid cracking and delamination.
To control crack defects, corresponding measures can be taken as follows: (1) Mixed
wires and optimization of their compositions; (2) Decrease the cooling rate during the
deposition process (3) Other measures to improve strength rather than solution treatment.
2.4 Current methods for quality improvement in WAAM process
Generally, WAAM parts require post-process treatment to improve material properties,
reduce surface roughness and porosity, and remove residual stress and distortions. By
CHAPTER 2 LITERATURE REVIEW
25
appropriate application of post process, the majority of issues that influence deposition
quality can be mitigated or eliminated. Presently, several post-process treatment technologies
have been reported to improve part quality in the WAAM process. This section will review
these post-process techniques, both their features and limitations.
2.4.1 Post heat treatment
Post-process heat treatment is widely used WAAM to reduce residual stress, enhance
material strength and as a method of hardness control. The selection of a suitable heat
treatment process depends on additive the target material, additive manufacturing methods,
working temperature and heat treatment conditions. If the heat treatment state is set
incorrectly, the probability of cracking will increase under mechanical loading, as the
combination of existing residual stress with load stress exceeds the material’s design
limitation.
As summarized above, after heat treatment, the mechanical strength of
WAAM-fabricated parts improved significantly, with increase of 4%, 78%, 5% and 17%
being reported for titanium alloy, aluminum alloy, Nickel-based superalloys and intermetallic
Ti/Al alloy, respectively. In addition, post-process heat treatment plays an important role in
grain refinement, especially for WAAM-fabricated aluminum and Inconel alloy [67].
The decision to use post-process heat treatment depends on the material alloying system
and also the pre-heat treatment state. Generally, high carbon content materials must be heat
treated, while a few materials can be damaged by this technique. Therefore, the utilization of
post heat treatment process to WAAM component needs to consider the specific material and
its application.
CHAPTER 2 LITERATURE REVIEW
26
2.4.2 Interpass cold rolling
Rolling of the weld bead between each deposited layer has been shown to reduce
residual stresses and distortion [114]. Interpass cold rolling not only lowers residual stress,
but also brings more homogeneous material properties. In the WAAM process, the thermal
gradient with deposition layers and alternate re-heating and re-cooling process result in the
target part having anisotropic microstructural evolution and mechanical properties. The cold
rolling technique significantly reduces microstructural anisotropy through plastically
deforming the deposition. Figure 2-6 shows the schematic diagram of an interpass cold
rolling system developed at Cranfield University. A slotted roller is used o refine the
microstructure and enhance tensile strength in the longitudinal direction by supporting
external force [115, 116]. As shown in Table 2-8, both ultimate tensile strength and yield
strength in the build direction were improved through interpass cold rolling, which
contributes to more homogeneous material properties in the target component.
Interpass cold rolling also can play a critical role in the healing of hydrogen porosity in
WAAM-fabricated aluminum parts. Generally, high dislocation density is produced by the
rolling process. These dislocation can act as preferential sites for atomic hydrogen absorption
[117] and as well as ‘pipes’ for the hydrogen, allowing to diffuse to the surface [118]. Table
2-9 summarizes the outcomes documented in the literature, in terms of the pore incidence and
size distribution in aluminum components fabricated using WAAM with interpass cold
rolling. The porosities existing in as-fabricated component can be reduced or even eliminated
when interpass cold rolling is applied [75].
This technique is only feasible for simple deposited parts, such as straight walls, due to
CHAPTER 2 LITERATURE REVIEW
27
the geometrical limitation of the rolling process. For more complex components with curves
and corners, special flexible tooling need to be developed to achieve an effective rolling
process, thus limiting the range of industrial application. Cold rolling technique will also
reduce residual stress, but the ability to reduce overall part distortion is yet to be proven [26,
119].
Figure 2-6 Schematic diagram of WAAM with cold rolling process [104]
Table 2-8 Material properties of component fabricated using WAAM with interpass cold
rolling
Materials Condition Microstructure YS[MPa] UTS[MPa] EL[%] Hardness[HV] Ref.
Ti6Al4V
AF Widmanstätten α + Martensite 807 ± 2a
853 ± 10b
903 ± 3a
945 ± 6b
22.8 ± 0.2a
13.8 ± 1.2b /
[114,
116,
120,
121]
RL:50 Grain size decrease from
average 200μm to 100μm
916 ± 22a
911 ± 11b
1022 ± 3a
1006 ± 5b
14.2 ± 1.3a
10.0 ± 0.5b /
RL:75 993 ± 31a
1028 ± 5b
1078 ± 3a
1077 ± 18b
13.0 ± 1.0a
13.1 ± 1.8b /
Al 2319
AF
/
132 260 15.5 /
[120] RL:15 145 274 13 /
RL:30 190 288 10.3 /
RL:45 244 304 7.3 /
AISI
301LN
AF Austenite phase 292 697 54 220
[115]
CR:20
transformed to α'-rnartensite
+ deformed austenite
750 1034 23 345
CR:40 1239 1321 6 456
CR:60 1529 1062 3 522
CR:80 1900 1925 2 610
Al-6.3Cu AF Fine dendrites + equiaxed grains 128 ± 2a
133 ± 5b
262 ± 4a
264 ± 2b
15.8 ± 0.3a
18.6 ± 1.5b 68.3 [44]
CHAPTER 2 LITERATURE REVIEW
28
RL:15
With the increase load, the
grain size is decreased with
freeform deformation
142 ± 4a
146 ± 3b
270 ± 7a
271 ± 9b
14.6 ± 0.4a
15 ± 1.4b 77.4
RL:30 188 ± 4a
195 ± 2b
286 ± 3a
293 ± 6b
12 ± 0.6a
13.2 ± 0.2b 90.1
RL:45 241 ± 2a
249 ± 1b
312 ± 8a
323 ± 9b
7.4 ± 0.6a
8.6 ± 0.4b 102.3
AF: as-fabricated; RL: Rolling loads (KN); CR: Cold reduction (%);
a: In build direction, b: Orthogonal to build direction.
Table 2-9 The distribution of porosity in aluminum part using WAAM with cold rolling.
Materials Condition Layers Height
(mm)
Deformation
(%)
Number of pores
(in area of 120 mm2 )
Pores diameter
(mm) Reported by
Al 2319
AF 21 49.4 / 614 13.5
Gu et al.[76]
RL:15 25 50.6 13.9 192 12.5
RL:30 30 49.4 30.0 5 8.8
RL:45 38 49.8 44.2 Elimination Elimination
Al 5087
AF 30 49.5 / 454 25.1
RL:15 35 49.6 14.1 336 13
RL:30 45 51 31.3 11 9.6
RL:45 55 49.5 45.4 Elimination Elimination
AF: as-fabricated; RL: Rolling loads (KN);
2.4.3 Interpass cooling
Recently, interpass cooling has been developed and evaluated at the University of
Wollongong. Figure 2-7 presents the schematic diagram of a WAAM system with interpass
cooling. The moveable gas nozzle, which supplies argon, nitrogen or CO2 gas, is used to
provide active, or forced, cooling on the fabricated part during and/or after deposition of each
layer. Using such rapid cooling, the in-situ layer temperature and heat cycle can be controlled
within a range to obtain the desired microstructure and mechanical properties. This process
may also potentially reduce residual stress and distortion, although this aspect has not been
investigated.
An initial feasibility study shows promising results when using forced interpass cooling
CHAPTER 2 LITERATURE REVIEW
29
with compressed CO2 to fabricate Ti6Al4V thin-walled structures, as shown in Figure 2-8. It
was found that interpass cooling produces less surface oxidation, refined microstructure,
improved hardness and enhanced strength. In addition, manufacturing efficiency is
significantly improved due to the reduction of dwell time between deposited layers. More
detailed research findings will be presented in future.
Figure 2-7 Schematic diagram of the combined WAAM gas cooling process
(a)
CHAPTER 2 LITERATURE REVIEW
30
(b)
Figure 2-8 Mechanical properties of Ti6Al4V part produced by WAAM with interpass
cooling using CO2 gas:(a) Hardness; (b) Tensile strength and elongation
2.4.4 Peening and ultrasonic impact treatment
Peening and ultrasonic impact treatments (UIT) have been used in welding applications
to reduce local residual stress and improve weld mechanical properties. Both techniques are
cold mechanical treatments that impact the weld surface using high energy media to release
tensile stress by imposing compressive stress at the treatment surface. Usually, the
mechanical peening process produces compressive stresses at a limited depth below the
component surface, such as around 1–2mm in carbon steels [122]. Ultrasonic impact
treatment produces grain refinement and randomizes orientation, thus contributing to
mechanical strength improvement. It is reported that after ultrasonic impact treatment, the
surface residual stress of WAAM-fabricated Ti6Al4V part can be reduced to 58% and the
microhardness can be increased by 28% compared to the as-fabricated sample. Also, the
surface-modified layers undergo plastic deformation with significant grain refinement and
dense dislocations [123]. The ultrasonic impact treatment is limited by penetration depth,
CHAPTER 2 LITERATURE REVIEW
31
which is up to 60 μm below surface. Therefore, although both techniques are good
post-process treatments, they have negligible effect on internal residual stresses of large metal
part fabricated using WAAM.
2.5 Summary and scope of this work
A review of recent technology development in WAAM process is presented with a focus
on the microstructure, mechanical properties, process defects and post-process treatment. In
combination of materials characteristics and WAAM techniques features, a quality-based
framework will be proposed in final section of this thesis, aiming to achieve high-quality and
defect-free WAAM components.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
32
Chapter 3 EXPERIMENTAL INSTRUMENTS
AND METHODOLOGIES
In this chapter, the experimental instruments and methodologies used in the thesis will
be briefly illustrated, and the details, such as experimental parameters, will be displayed in
the following chapters.
3.1 Materials
The materials used in this study were commercial ASTM B863 grade 5 Ti-6Al-4V alloy
wire with a diameter of 1.2mm and ASTM B265 grade 5 Ti-6Al-4V sheet with dimensions of
200 mm × 150 mm × 6 mm. The chemical composition is listed in Table 3-1.
Table 3-1 Chemical composition of Ti-6Al-4V (wt.%).
Composition Al V C Fe H N O Ti
Wire (ASTM B863)[124] 6.20 4.0 0.08 0.40 0.015 0.05 0.20 Bal.
Substrate (ASTM B265)[125] 6.10 4.0 0.08 0.30 0.015 0.03 0.20 Bal.
3.2 Experimental instruments
The apparatus used in this thesis mainly consists of a 200A-rated GTAW power source,
“cold” wire feeder, water cooling unit and travel mechanism, as shown in Figure 3-1. A 150
mm long trailing gas shield was placed behind the GTAW torch to provide localized
protection against oxidation during the cooling process.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
33
(a) (b)
(c)
Figure 3-1 The apparatus used in this work: (a) 200A-rated GTAW power source; (b) wire
feeder; (c) travel mechanism and trailing gas shield
3.3 Process monitoring
During manufacturing process, a high speed camera designed to record the dynamic arc
and the metal transfer process at a frame rate of 1 kHZ, as shown in Figure 3-2. the AF-S
NIKKOR 70-200mm*1:4 ED VR camera lense was used. The experimental data obtained
from the high speed camera profiler was processed using MATLAB.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
34
Figure 3-2 High speed camera used in this study
An infrared (IR) pyrometer system with an optical fiber was used to measure the
temperature of the substrate and the deposited part during manufacturing, as shown in Figure
3-3. The temperature of the pyrometer system was read by the developed LabView software.
Figure 3-3 IR pyrometer used in this study
A 3D laser profile scanner with a resolution of 0.02 mm was adopted to measure the
deformation of samples, as shown in Figure 3-3. To improve measurement accuracy which
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
35
involves random errors, each sample profile was scanned 200 times. The experimental data
obtained from the laser profiler was processed using MATLAB.
Figure 3-4 3D laser scanner used in this work
3.4 Metallography
After sample preparation, microstructural characterization of the specimens was carried
out using stereo microscopy, optical microscopy (OM) and scanning electron microscopy
(SEM). Chemical composition of precipitates was analysed by energy dispersive
spectroscopy (EDS) detectors, which are attached to SEM.
3.4.1 Sample preparation for metallography
The specimens for microstructural observation (OM and SEM) were cut using a Wire
Electrical Discharge Machining. Afterwards, these specimens were hot mounted with a
CitoPress-20 hot mounting press (Figure 3-5), and then polished in a standard metallographic
way by Struers automatic polisher Tegrapol 21 (Figure 3-6). The detailed grinding and
polishing procedures are provided in Table 3-2. After that, the specimens were etched in
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
36
Kroll’s reagent solutions that consist of 100ml distilled water (H2O), 5ml concentrated nitric
acid (HNO3) and 3ml 3% hydrofluoric acid (HF).
Figure 3-5 Struers CitoPress-20
Figure 3-6 Struers Tegrapol-21
Table 3-2 Sample preparation procedures for microstructural analysis.
Procedure Surface Force, N Time, minute Solution
Grinding 1200# SiC paper 25 2 Water
Polishing 15μm MD-Pan cloth 25 15 Water-based lubricant
Polishing 0.25μm MD-Chem cloth 25 10 50% OPS
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
37
3.4.2 Stereo microscopy and optical microcopy (OM)
Based on visible light and a series of lenses, a Leica M205A stereo microscopy (Figure
3-7) was used to acquire images of cross-section of specimens and a Leica DMR optical
microcopy (Figure 3-8) was used to magnify images. The horizontal and vertical section is
for microstructure observation, including precipitation and grain size. The grain size was
measured by the software of Image 2 using planimetric method.
Figure 3-7 Leica M205A stereo microscopy
Figure 3-8 Leica DMR optical microscopy
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
38
3.4.3 Scanning electron microscopy (SEM)
SEM is a type of electron microscope through scanning the surface of a sample with a
focused beam of electrons. In this thesis, a JEOL 7500 operating SEM equipped with an
energy dispersive X-ray spectrometer (EDS) was used to observe the fracture surfaces of
tensile specimens, as shown in Figure 3-9. Atomic concentrations of precipitates were
analysed by EDS if special composition of phase on fracture surfaces of specimens needs to
be further confirmed.
Figure 3-9 JEOL JSM-7500 SEM
3.4.4 X-ray diffraction (XRD)
X-ray diffraction (XRD) was used to determine the phase constitution of specimens. In
this thesis, XRD testing was performed on a GBC MMA X-ray diffractometer using
monochromatic Cu K radiation (wavelength = 1.5418Å) at an accelerating voltage of 35
kV and a current of 28.6 mA with a scanning step of 0.02º and a scanning speed of 4 º/min in
the range of 30º ~ 90º, as shown in Figure 3-10.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
39
Figure 3-10 GBC MMA X-ray diffractometer
3.5 Mechanical properties
Mechanical properties of the specimens were measured, focusing on tensile strength and
hardness testing and aiming at evaluating the strength, hardness and ductility properties under
different WAAM process.
3.5.1 Tensile testing
In this thesis, tensile tests were conducted using a MTS 370 (Figure 3-11) at room
temperature. The ultimate tensile strength (UTS), yield strength (YS) and elongation were
measured during test. The diagram of machined specimens for tensile testing is shown in
Figure 3-12, and the detailed dimension data is listed in Table 3-3.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
40
Figure 3-11 MTS370 tensile machine
Figure 3-12 The diagram of machined tensile specimen.
Table 3-3 Dimensions for tensile specimen.
Standard specimens, Sheet-type Dimension
G-gage length 10.0 ± 0.10 mm
W-Width 2.0 ± 0.25 mm
T-Thickness 1.5 mm
R-Radius of fillet, min 2.0 mm
L-Over-all length, min 34.0 mm
A-Length of reduced section, min 2.0 mm
B-Length of grip section, min 10.0 mm
C-Width of grip section, approximate 8.0 mm
3.5.2 Hardness testing
The Vickers hardness testing was performed on a Via-F Vickers innovative automatic
tester (Figure 3-13) with a load of 0. 98 N and a dwell time of 10 s according to ASTM:
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
41
F2924 standard. Specimen preparation should be performed in accordance with standard
guide for preparation of metallographic specimens. The corresponding results were presented
graphically.
Figure 3-12 The Vickers innovative automatic testers
3.6 Corrosion resistance
The electrochemical corrosion tests were conducted on a CHI 600C electrochemical
working station (manufactured by CH Instruments, Inc.) with a conventional three-electrode
cell in 3.5 wt.% NaCl solution at room temperature. A platinum sheet and saturated calomel
electrode (SCE) were used as counter electrode and reference, respectively, as shown in
Figure 3-14.
CHAPTER 3 EXPERIMETN INSTRUMENTS AND METHODOLOGIES
42
Figure 3-13 Electrochemical workstation
3.7 Summary
This chapter insulates the commonly used experimental instruments and methodologies
in the thesis. The particular and detailed experimental procedures as well as experimental
results will be presented and discussed in the relative chapters
CHAPTER 4 PROCESS STABILITY
43
Chapter 4 PROCESS STABILITY
4.1 Introduction
The heat accumulation is a critical factor that influences the stability of the WAAM
process in terms of the geometrical accuracy, deposition defects, microstructural evolution
and material properties of as-fabricated parts. Ma et al. [92] carried out an experiment in
which a simple Ti-Al part was fabricated using GT-WAAM. It was found that the alpha phase
fraction in the microstructure decreases by nearly threefold when the interpass temperature
changes from 100 to 500 , which results in a decrease in hardness values. Shen et al.
[126]investigated the fabrication of Fe-Al materials using a similar WAAM process and
reported that poorly controlled interpass temperature is likely to produce longitudinal
cracking and high residual stress in the first few layers. In addition, considerable research has
focused on arc stability and metal transfer behaviour of the WAAM process. Wang et al.
[86]claims that during the arc-wire deposition process, the distance between the trailing end
and center of molten pool was increased by 1.95 mm from 1st layer to 5th layer due to the
increased heat accumulation. Similar results could also be found in earlier literature regarding
the analysis of thermal behaviours during multi-layer rapid prototyping fabrication[127].
Another study conducted by Denlinger et al. [43] found that the distortion and residual
stresses in as-fabricated titanium and nickel alloy parts were significantly affected by the
inter-layer dwell time that is directly related to the thermal characteristics. Zhou et al.
[128]developed a three-dimensional model to simulate the arc shape and metal transfer
behaviours occurring in the WAAM process. The distribution of thermal conductivities and
CHAPTER 4 PROCESS STABILITY
44
molten pool characteristics for single-bead as well as overlap deposition were investigated.
Their results show that the high temperature region of the molten weld pool for overlapping
deposition is narrower than that of single-bead deposition, due to the smaller net heat flux of
overlapping deposition. Although these simulation and experimental studies have provided
some useful information, the underlying mechanisms of arc characteristics and metal transfer
behavior associated with heat accumulation are still poorly understood due to the complexity
of the WAAM process.
This section investigates the influence of heat accumulation on process stability during
the fabrication of Ti6Al4V parts using GT-WAAM with localized gas shielding. The
difference in temperature variations between the substrate and in-situ layers is discussed. The
results provide insight into the way in which the interpass temperature is measured.
Furthermore, the surface morphology, geometrical features, arc characteristics and metal
transfer behaviours in different layers during deposition are compared to identify variation
trend associated with heat accumulation. Although the research outcomes are not a direct
physical explanation of the thermal mechanism, this section offers better understanding of the
effects of heat accumulation on process stability of the wire arc additive manufacturing
process, which will benefit further process optimization and control.
4.2 Experimental Procedures
4.2.1 Experimental setup
The GT-WAAM system used in this section, as shown in Figure 4-1. The wire filler was
fed into the front of the weld pool at an angle of 60°to the GTAW torch, which produces a
CHAPTER 4 PROCESS STABILITY
45
stable metal transfer condition. Welding grade argon (99.995% purity) is used for both torch
and trailing shield. A build wall was produced by depositing 15 passes on the as-received
substrate plate. The main process parameters are listed in Table 4-1.
Figure 4-1 Schematic illustration of GT-WAAM process
Table 4-1 Process parameters for GT-WAAW
Process parameters Details
Deposition current 110 A
Arc voltage 12 V
Travel speed 95 mm/min
Wire feed speed 1000 mm/min
Distance between the electrode and workpiece 3 mm
Angle between the electrode and the filler wire 60°
Flow rate of argon in GTAW torch 10 L/ min
Flow rate of argon in trailing shield 10 L/ min
Post flow shielding time 30 seconds
Dwell time between layers 180 seconds
4.2.2 Measurement of interpass temperature
Thermocouples and an IR pyrometer were used to measure the temperature of the
substrate and the deposited part, respectively. Figure 4-2a shows the operating principle of
CHAPTER 4 PROCESS STABILITY
46
the measurement system. Two thermocouples were fixed on the substrate for monitoring the
substrate temperature, and the IR pyrometer was positioned to measure the in-situ layer
temperature immediately before deposition of the subsequent layer. The specific
measurement locations are shown in Figure 4-2b.
(a)
(b)
Figure 4-2 Schematic diagram of the temperature measurement system: (a) overall system (b)
measurement locations for the pyrometer and thermocouples (TC)
The material emissivity was calibrated before using the infrared pyrometer since the
emissivity of the material varies with bulk, surface roughness, oxidization, and especially
with surface temperature, as reported by Hagqvist et al.[129]. The temperature (TIR) from a
CHAPTER 4 PROCESS STABILITY
47
signal-wavelength IR pyrometer is expressed as[130]:
𝑇𝐼𝑅 =𝐶𝑇
𝐶−𝜆𝑇𝑙𝑛( 𝐹𝛾) (4-1)
where T is the nominal absolute temperature, ε is the emissivity of the material surface, F is
the ratio of the capturing field of view of pyrometer, γ is the attenuation caused by the media
(air) from sample to the pyrometer sensor, λ is the wavelength of the pyrometer, and C is the
second radiation constant(C=1.4387 × 104 μm∙K). For a small emissivity error, Δε, the
following expression describes the magnitude of the temperature error [130]:
∆𝑇 =𝜆𝑇2
𝐶∙
∆ (4-2)
A number of calibration tests were carried out on a single deposited Ti6Al4V bead
(smooth silver surface) that was continuously heated from 50 to 500 , followed by
natural cooling process back to room temperature, as shown in Figure 4-3. Two
thermocouples were attached on the bead surface as the temperature reference for calibration.
As no shielding gas was used during this calibration, slight discoloration due to oxidation on
the Ti6Al4V bead surface was observed, which provides an appropriate simulation for the
actual heating and cooling cycle during actual deposition. Through multiple tests, a constant
emissivity of 0.45 was determined to minimize the temperature measurement errors due to
small variation in emissivity compensation. Figure 4-3 shows the temperature calibration
curves from both pyrometer and thermocouples (actual temperature) with emissivity value set
as 0.45. It can be seen that the relationship between temperature acquired from pyrometer and
actual temperature can be expressed as the following 2nd order polynomial relationship:
y = −0.0007𝑥2 + 1.3223𝑥 − 16.07 (4-3)
CHAPTER 4 PROCESS STABILITY
48
where x is the temperature acquired from pyrometer using a constant emissivity of 0.45, and y
is the actual temperature of the metal measured with thermocouples. Using the calibration
formula from Eq.3, it was found that the largest temperature measurement error from 50 to
500 is less than 10 .
Figure 4-3 Pyrometer calibration curve for Ti6Al4V using fixed emissivity of 0.45
4.3 Results and discussion
4.3.1 Interpass temperature and heat accumulation
Figure 4-4 shows the variation in temperature at both substrate and in-situ layer
immediately before each subsequent deposition pass. It can be seen that the average substrate
temperature experiences a rapid increase during the first five passes, and then reaches an
equilibrium value. In contrast, the interpass temperature continues to climb up to 15 passes.
Previous researchers have used the substrate temperature as a key variable for investigating
the thermal behaviour of additive manufacturing processes, as claimed by Ding et al. [33] and
CHAPTER 4 PROCESS STABILITY
49
Denlinger et al.[43] However, based on the experimental outcomes from this study, it is
evident that there is a large discrepancy between the temperature measured at the substrate
and the actual interpass temperature, particular if the dwell time between layers is short. If the
temperature measured at substrate is taken to be the interpass temperature, it will cause large
errors. Direct measurement of the layer surface temperature using noncontact IR pyrometer
provides far more accurate and reliable data.
Figure 4-4 The variation in temperature and heat accumulation during fabrication
In order to estimate heat accumulation quantitatively, the equation of Specific Heat
Capacity is utilized to calculate the heat accumulation at each layer and is expressed as:
𝑄𝑗 = 𝐶 ∙ ∑ 𝑚𝑗𝑛𝑗=1 ∙ ∆𝑇 = 𝐶 ∙ ∑ 𝑚𝑗
𝑛𝑗=1 ∙ (𝑇𝑗 − 𝑇1) (4-4)
𝑚𝑗 = 𝜌 ∙ 𝜋𝑟3 ∙ 𝑣 ∙ 𝑡𝑗 ∙ 𝜀 (4-5)
where C is the heat capacity of Ti6Al4V, approximately 0.619J/(g·K), m is the mass of jth
deposited layer, Tj is the interpass temperature of jth layer, 𝜌 is the density of Ti6Al4V, r is
CHAPTER 4 PROCESS STABILITY
50
the diameter of the welding wire, v is the wire feeding speed, tj is the wire feeding time for jth
layer, and 𝜀 is the material efficiency, which is between 0.95 to 0.98 for the GT-WAAM
process. During fabrication, the difference in interpass temperature between subsequent
layers determines the final heat accumulation of the as-fabricated part. The detailed data for
each layer are listed in Table 4-2 and the heat accumulation results are presented in Figure
4-4. The heat accumulation increased almost linearly with the number of layers as expected.
Table 4-2 Detailed process data for heat accumulation calculation.
i 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15
t/s 95 94 97 98 98 97 98 99 98 95 94 95 93 93 90
m/g 7.57 7.49 7.73 7.81 7.81 7.73 7.81 7.88 7.81 7.57 7.49 7.57 7.41 7.41 7.17
T/ 139 165 194 204 209 215 218 236 241 250 258 266 287 292 296
Q/KJ 1.85 3.12 5.16 7.14 9.10 11.13 13.13 15.80 18.04 20.49 22.94 25.49 29.00 31.51 33.99
4.3.2 Bead appearance and geometrical features
Figure 4-5 compares the appearance of the first layer with that of the top layer. The
surface of the first layer shows a clean silver surface with no sign of oxidation (Figure 4-5a),
whereas obvious oxidation can be observed at the top layer. This phenomenon indicates that
interpass temperature has increased with increasing build height due to the heat
accumulation.
CHAPTER 4 PROCESS STABILITY
51
Figure 4-5 The bead appearance of (a) the first layer, (b) the top layer
Hong and Shin [131] pointed out that the appearance of Ti6Al4V alloy can be used for
visual assessment of surface imperfections, degree of oxidation as well as the effectiveness of
inert gas shielding against atmospheric contamination. As for the surface discoloration, the
surface color changes from silver, light straw through dark straw, light blue, dark blue, to
grey and powdery, in order of increasing oxidation and contamination. As shown in Figure
4-5, the surface colors for the base layer and the top layer are silver and dark blue or gray,
respectively, which indicates heavy oxidation on the surface of the top layer. Wang et al. [132]
mentioned that surface oxidation was sensitive to temperature and oxygen concentrations in
the atmosphere. With an increasing number of deposited layers during GT-WAAM of
Ti6Al4V, the surface of freshly deposited layers are exposed to increasingly higher
interpass temperatures with a constant interlayer dwell time due to the heat accumulation
effects, which produces increasingly severe oxidation. Hence, where a Ti6Al4V part is
fabricated using localized gas shielding in an otherwise open atmosphere, it is necessary to
control the interpass temperature to ensure that surface oxidation of the most recent deposit is
kept at an acceptably low level, to avoid further contamination of the next deposited layer.
Figure 4-6 shows the width of the cross section of the as-fabricated wall along the build
CHAPTER 4 PROCESS STABILITY
52
height, which experiences a sharp increase from the 1st layer to the 5th layer. In the first few
layers, heat is readily conducted to the substrate, leading to a relatively faster cooling rate and
narrower bead. As the wall is built up, the conductive thermal resistance to the substrate is
significantly increased. Titanium alloys exhibit poor thermal conductivity, so a decreasing
fraction of the heat input from the welding process is conducted away through the substrate
(which is now at an elevated temperature). An increasing amount of heat is dissipated to the
surrounding atmosphere via convection and radiation, but these temperature-reducing
mechanisms are less effective than direct conduction to a cool substrate, which leads to the
slower heat dissipation condition of the molten pool and wider bead at higher layers. As more
layers are deposited with a fixed dwell time, the deposition width of the wall becomes steady
as the heat input and dissipation reach a balance.
Figure 4-6 Width of the build cross-section along the building height
CHAPTER 4 PROCESS STABILITY
53
4.3.3 Stability of weld pool and arc behaviour
4.3.3.1 Arc behaviours
Arc profiles acquired from the high speed camera system were processed using
MATLAB, and the arc length measured at each layer is presented in Figure 4-7. It can be
seen that the average arc length of the first few layers increases with the increasing number of
deposition layers, then it tends to be stable, which is consistent with the change in the bead
geometry as shown in Figure 4-6. The error bar in the Figure 4-7, represents the
measurement errors of arc length using a total of 500 photographs for each layer.
Figure 4-7 The arc shape evolution in GT-WAAM of Ti6Al4V
In general, the arc length is determined by the distance between tungsten tip and molten
pool surface. For the GT-WAAM process using localized shielding, even though slight
oxidation occurs at the upper layers, the upper surface is still relatively flat and will not be the
cause of an inconsistent arc length. However, an obvious expansion of arc profile in the first
few layers was observed, as shown in Figure 4-7, which indicates that the distance between
CHAPTER 4 PROCESS STABILITY
54
tungsten tip and molten pool increases significantly due to the considerable change in bead
geometry. According to Wahab et al.[133], the mathematical model of heat transfer in the
molten pool is expressed in the following form:
𝑘𝛿2𝑇
𝛿𝑥2 + 𝑘𝛿2𝑇
𝛿𝑦2 + 𝑘𝛿2𝑇
𝛿𝑧2 + 𝑘𝛿𝑄
𝛿𝑡2 = 𝜌𝐶𝛿𝑇
𝛿𝑡 (4-6)
With the boundary conditions represented as:
𝑘𝛿𝑇
𝛿𝑛+ ℎ(𝑇 − 𝑇0) + 𝜎𝜀(𝑇4 − 𝑇0
4) = 0 (4-7)
Based on the assumption of a quasi-steady state approximation, equation (4-7) can be
rearranged as follow:
𝑘𝛿2𝑇
𝛿𝑥2 + 𝑘𝛿2𝑇
𝛿𝑦2 + 𝑘𝛿2𝑇
𝛿𝑧2 − 𝜇𝛿𝑄
𝛿𝑥= −𝜇𝜌𝐶
𝛿𝑇
𝛿𝑥 (4-8)
where Q is the internal heat energy released or consumed per unit volume (J∙mm-3), T is the
temperature related to the heat accumulation, T0 is the ambient temperature, t is time, k is
thermal conductivity ( W∙(mm∙°C)-1 ), r is density, C is specific heat of the material (J∙
(g∙°C)-1), h is a convection coefficient, σ is the Stefan Boltzman constant and ε is emissivity.
As presented by numerical simulation from these conductive models, more heat energy from
the arc and molten pool can increase molten pool length and width, which has also been
reported by Wang et al.[86]. According to the law of mass conservation, the height of molten
pool will decrease in order to make up for the increase in its width and length. In this study,
the arc length increased with the variation in bead geometry at first few layers. After five
layers’ deposition, it stabilizes with further deposition due to constant bead geometry
associated with heat equilibrium, as shown in Figure 4-8. This is in good agreement with the
CHAPTER 4 PROCESS STABILITY
55
change in bead width along the building height as depicted in Figure 4-7.
Figure 4-8 The schematic diagram of the changes in arc length and bead height for different
layers: (a) layer 1, (b) layer 5, (c) layer 10 and (d) layer 15
In addition, according to findings from Zhou et al.[128], metal vapor produced from the
molten pool tends to increase the heat and radiation loss in low temperature areas of the arc,
leading to potential arc constriction. To be specific, the isotherms shrinking toward the arc
center implies that intensified metal vapor could induce changes in the arc behavior at
different layers due to accumulated heat serving as a preheating heat source, which
contributes to the melting process of filler and base layer[134]. Therefore, the variation in arc
shape can be explained through the combined effects of deposited layer geometry and arc
constriction although it has no direct link with heat accumulation in the additive
manufacturing process.
CHAPTER 4 PROCESS STABILITY
56
4.3.3.2 Metal transfer behaviours
As mentioned above, with increasing built height, the arc shape expands slightly in first
few layers and then becomes nearly constant. This phenomenon also has certain impact on
the metal transfer behaviour, which has a direct influence on the bead formation and
geometrical accuracy of the deposited metal. In terms of cold-wire GTAW, the metal transfer
modes can be mainly categorised into four types: intermittent wire melting, uninterrupted
bridging transfer, interrupted bridging transfer and free flight transfer[135]. For the
GT-WAAM process, the uninterrupted bridging transfer mode forms a smooth and consistent
layer appearance, while the irregular free flight mode tends to produce humps and hollows on
the surface, as reported by Geng et al.[41].
Figure 4-9 tracks the specific metal transfer process at different layers. It can be seen
that the metal transfer state changes with increasing number of layers, especially in the first
few layers. At layer 1 where deposition is on the substrate, the metal transfer is through
uninterrupted bridging transfer mode, while it gradually changes into interrupted bridging
mode in layer 5. Figure 4-10 provides more detailed images of the metal transfer at specific
stages during the deposition. It can be observed that the width of liquid bridging decreased
gradually from 1st to 5th layer, which indicates that the distance between the tungsten tip and
molten pool has increased due to the change in bead geometry. Furthermore, the free flight
transfer mode was occasionally observed in the upper layers due to some distortion of the
part under the successive thermal cycles.
CHAPTER 4 PROCESS STABILITY
57
(a)
(b)
(c)
CHAPTER 4 PROCESS STABILITY
58
(d)
Figure 4-9 Metal transfer process at: (a) layer 1, (b) layer 5, (c) layer 15; (d) Interrupted
transfer occurring from layer 5 to 15
Figure 4-10 Metal transfer behaviour at different layers: (a) layer 1; (b) layer 5; (c) layer 15
CHAPTER 4 PROCESS STABILITY
59
It is worth noting that the metal transfer frequency has also experienced a slight increase
due to the variations in arc profiles, as shown in Figure 4-9. Considering that the variation in
the arc shape and molten pool geometry occurs, the forces acting on the droplet also change
significantly during the build-up. According to the static force balance theory introduced by
Waszink and Graat [136], the major forces acting on the droplet include gravity (Fg),
electromagnetic force (Fem), surface tension (Fσ) vapor jet force (Fv), and plasma drag force
(Fa). For the GTAW-AM process, the balance of these forces determines the metal transfer
process, as described in Figure 4-11. The detaching force (Fd) is defined as the resultant force
on the droplet in the vertical direction, which can be derived by the following
expression[137]:
𝐹𝑑 = (𝐹𝑔 + 𝐹𝑒𝑚 + 𝐹𝑎 + 𝐹𝑝𝑐𝑜𝑠60°) − (𝐹𝑣 + 𝐹𝜎𝑆𝑖𝑛30°)
=4
3𝜋𝑟𝑑
3𝜌𝑑𝑔 +𝜇0𝐼2
4𝜋(
1
2+ 𝑙𝑛
𝑟𝑢
𝑟𝑖) + 𝐹𝑝𝑐𝑜𝑠60° + 𝐶𝑑𝐴𝑝 (
𝜌𝑓𝑣𝑓2
2) (4-9)
−𝑚0
𝑑𝑓𝐿𝐽 − 2𝜋𝑅𝜎𝑠𝑖𝑛30°
where 𝐹𝑝 is the driving force of feed wire, rd is the droplet radius, ρd is the density of droplet,
g is the gravitational constant, μ0 is the magnetic permittivity, I is the welding current, ri is the
exit radius of the current path, ru is the entry radius of the current path, which is equal to the
radius of anode area, Cd is the drag coefficient, Ap is the projected area on the plane
perpendicular to the fluid flow, ρf and vf are the density and fluid velocity of the plasma,
respectively, m0 is the total mass vaporized per second per ampere, I is the welding current, J
is the vapor density, R is the radius of the electrode, and σ is the surface tension coefficient.
As the constriction of the arc induces an increase in electromagnetic force (Fem) acting on the
CHAPTER 4 PROCESS STABILITY
60
droplet at the first few layers, the requirement of gravitation force (Fg) for the droplet is
reduced according to equation (9). Therefore, the droplet in the liquid bridging area would be
detached at a smaller size and the metal transfer frequency increased accordingly.
Figure 4-11 Forces acting on droplet during Ti6Al4V GT-WAAM process
4.4 Conclusion
In this section, Ti6Al4V alloy has been used as the build material for the GT-WAAM
process using localized gas shielding. Based on the in-situ measurement of the temperature at
each layer, the effects of heat accumulation on stability of deposition, oxidation, geometrical
shape, arc characteristics and metal transfer behaviour were investigated. The main
conclusions are drawn as follows:
(1) By using a noncontact IR pyrometer for direct temperature measurements at the
deposition location on the wall, more accurate and reliable information on interpass
temperature can be obtained when compared to measurements from thermocouples attached
at preset locations on the substrate, due to the existence of heat accumulation.
(2) Due to the changes in heat dissipation path along the deposition height and
CHAPTER 4 PROCESS STABILITY
61
decreasing cooling rates, the bead geometry varies in the first few layers, which leads to
slight changes in the distance between electrode and molten pool.
(3) The slight variation in arc shape and metal transfer behaviour along the building
height can be explained by the combined effects of both increased distance between electrode
and the surface of molten pool as well as arc constriction resulting from intensified metal
vapor, even though it is not directly related to the thermal behaviours in the WAAM process.
In summary, the influence of heat accumulation on the deposition stability is proven to
be very significant during the TiAl64V GT-WAAM process using localized shielding.
Therefore, to achieve better geometry and more stable metal transfer, interpass temperature
must be strictly controlled during the WAAM of Ti6Al4V.
CHAPTER 5 MATERIAL PROPERTIES
62
Chapter 5 MATERIAL PROPERTIES
5.1 Introduction
Complex thermal behaviour during fabrication plays an import role in material properties
of components manufactured using Wire Arc Additive Manufacturing (WAAM) technology.
It has been demonstrated that the WAAM process is capable of producing high-quality
Ti6Al4V components with comparable mechanical properties to those produced by
conventional manufacturing processes[26]. However, the control of thermal behaviour, which
influences microstructural evaluation, oxidation and defect generation, is a challenging task.
Recently, through numerical simulation, Vastola et al.[138] reported that heat can transfer
through several deposited layers to induce microstructural evolution during the electron beam
melting (EBM) additive manufacturing (AM) process. Romano et al. [139] investigated the
temperature distribution and geometry of deposited Ti6AL4V using both laser and EBM AM
processes. Their results show that the heat accumulated within the immediate area of the laser
beam will result in a wider melt pool. However, the temperature distribution of EBM is more
even because of the higher effective conductivity of the powder bed. Li and Gu [140]
observed that during the selective laser melting (SLM) process, heat accumulation can affect
the average temperature of the titanium powder bed, leading to a change of geometrical shape
and poor accuracy. Wu et al.[141] also found that the heat accumulation will affect the
geometry accuracy of the as built component in WAAM Ti6Al4V process, which indicates
the importance of in-process temperature monitoring and control. Although the effects of
thermal history, such as process temperature and cooling rate, on morphology and properties
CHAPTER 5 MATERIAL PROPERTIES
63
of materials have been studied by many researchers, little attention has been paid to the
influence of heat accumulation in gas tungsten wire arc additive manufacturing (GT-WAAM)
for Ti6Al4V alloy fabrication. It is believed that the microstructure of deposited part would
be influenced when the heat accumulates during manufacturing process, which eventually has
negative effects to the component property. A better understanding of the underlying
mechanism will be beneficial from the perspective of optimizing microstructural control.
In the present section, heat accumulation during GT-WAAM process using Ti6Al4V was
calculated by using in-situ interpass temperature measurements from an IR-pyrometer. The
surface oxidation, microstructure, grain size, crystalline phase and mechanical properties
from various locations of the deposited wall were investigated and discussed in terms of the
heat accumulation. Fracture experiments of the corresponding locations were also conducted
to reveal the correlation between microstructure and mechanical properties, and to assess the
possibilities of fabricating Ti6Al4V parts using WAAM with localized gas shielding.
5.2 Experimental Procedures
5.2.1 Experiment setup
A schematic diagram of experimental system used for this study is shown in Figure 5-1.
The GTAW process is locally shielded using welding-graded (99.995%) argon to prevent
oxidation of the deposited material. Both filler wire and substrate plate are Ti6Al4V. The
diameter of the filler wire is 1.2 mm and the dimension of the substrate plate is 200 mm ×
150 mm × 6 mm (length × width × thickness). The process parameters used for deposition are
provided in Table 5-1.
CHAPTER 5 MATERIAL PROPERTIES
64
Figure 5-1 Schematic diagram of the GT-WAAM system
Table 5-1 Process parameters for WAAM deposition
Process parameters Details
Deposition power Current 110A
Arc voltage 12V
Speed Travel speed 95mm/min
Wire feed speed 1000mm/min
Distance and angle Electrode to workpiece 3 mm
Electrode to filler wire 60°
Flow rate(argon) GTAW torch 10 L/ min
Trailing shield 10 L/ min
Time Post flow duration 35 seconds
Dwell time between layers 125 seconds
5.2.2 In-situ interpass temperature measurement
In order to accurately measure the in-situ process temperature, thermocouples and an IR
pyrometer were used, as shown in Figure 4-2. Two thermocouples were fixed on the
substrate for continuous monitoring of the substrate temperature, and an IR pyrometer was
CHAPTER 5 MATERIAL PROPERTIES
65
installed above the middle of deposited layer to measure the in-situ layer temperature
immediately after deposition.
Since the emissivity of the measured material is an essential measurement parameter, the
IR pyrometer needs to be calibrated in order to obtain an accurate temperature value. Several
calibration tests were carried out in the range of 50 to 500, and an emissivity of 0.45 was
established as the best fit to the temperature measurements from the thermocouples.
5.2.3 Heat accumulation calculation
The equation of Specific Heat Capacity Qj was used to calculate heat accumulation
during the WAAM process, as follows[141]:
𝑄𝑗 = 𝐶 ∙ ∑ 𝑚𝑗𝑛𝑗=1 ∙ ∆𝑇 = 𝐶 ∙ ∑ 𝑚𝑗
𝑛𝑗=1 ∙ (𝑇𝑗 − 𝑇1) (5-1)
𝑚𝑗 = 𝜌 ∙ 𝜋𝑟3 ∙ 𝑣 ∙ 𝑡𝑗 ∙ 𝜀 (5-2)
where C is the heat capacity of Ti6Al4V, approximately 0.619J/(g·K), m is the mass of jth
deposited layer, Tj is the interpass temperature of jth layer, 𝜌 is the density of Ti6Al4V, r is
the diameter of the filler, v is the wire feed speed, tj is the deposition time for jth layer, and 𝜀
is the material efficiency, which is between 0.95 to 0.98 .
Figure 5-2 shows the temperature changes in both substrate and upper-most deposited
layer during the fabrication process. It can be seen that after a few layers, the average
substrate temperature stabilizes, while the temperature climbs quickly in deposited layers
indicating the heat accumulation within the wall. Notably, the temperature difference between
the substrate and wall continues to increase after the fifth layer, and approaches a stable value
after thirteen layers. Based on equations (5-1) and (5-2), the detailed values of heat
accumulation are also presented in Figure 5-2 in blue.
CHAPTER 5 MATERIAL PROPERTIES
66
Figure 5-2The variation of temperature and heat accumulation in this study
5.2.4 Material characterization techniques
After grinding and polishing, the deposited specimens were etched in Kroll reagent (2%
vol HF, 6 vol% HNO3 with balance H2O) according to standard procedures for microstructure
observation using a Leica DMR optical microscope (OM). X-ray diffraction (XRD)
measurements were performed on a GBC MMA X-ray diffractometer using monochromatic
Cu Kα radiation at an accelerating voltage of 35 kV and a current of 28.6 mA with a scanning
step of 0.02º and a scanning speed of 4 º/min in the range of 30º ~ 90º. Tensile tests of
rectangular specimens (sample gauge length: 12 mm; width: 2 mm) were carried out on a
MTS370 load unit at a crosshead speed of 0.4 mm/min at room temperature [24]. The
fracture surfaces of tensile specimens were analyzed using a JEOL JSM-6490LA Scanning
Electron Microscopy (SEM) with an Energy Dispersive Spectrometer (EDS) at 20KV.
CHAPTER 5 MATERIAL PROPERTIES
67
5.3 Results and Discussion
5.3.1 Microstructure evolution
It is well known that for alpha+beta Ti6Al4V alloy, the thermal history determines the
formation of phases structures, including primary alpha, lathlike alpha, colony alpha and hcp
martensite alpha (alpha’), grain boundary alpha, acicular alpha and prior beta phases [73].
Figure 5-3 illustrates the microstructural distributions in the cross-section of the as-built
specimen and Figure 5-4 displays the detailed optical micrographs at the marked locations.
Due to the high cooling rate for the first few layers from direct contact with substrate, the
martensite alpha composed of long orthogonally oriented martensitic plates was formed into
lathlike matrix structures[142], as shown in Figure 5-4a. In general, this lathlike alpha tend
to grow in a long and narrow shape in the direction perpendicular to the liquid/solid interface
driven by the maximum temperature gradient during the solidification process[73]. As further
layers are deposited, more heat is accumulated in the wall and the process cooling rate
continue to reduce so that a fully lamellar alpha morphology is preferentially formed,
interwoven with basketweave structures (Figure 5-4b and c)[143]. At the top layers, large
colony alpha, which is decorated within prior beta grains and grain boundary alpha phase
(Figure 5-4 d, e and f), are formed. With the increase of heat accumulation, the process
temperature exceeds beta transus temperature Tβ (995°C[144]) at the uppermost layers. In
combination with a slow cooling rate, this results in the coarse colony alpha structures [145].
It is worth noting that all regions on the fabricated wall have finer grains than that of the base
metal (BM) (Figure 5-4 b), which is reflected in differences of the mechanical properties of
these regions.
CHAPTER 5 MATERIAL PROPERTIES
68
Figure 5-3 Cross-sectional macrograph of the deposited wall.
(a) (b)
CHAPTER 5 MATERIAL PROPERTIES
69
(c) (d)
(e) (f)
Figure 5-4 Optical micrographs of the corresponding regions: a, b, c, d, e, f in Figure 5-3
Figure 5-5 shows the optical micrographs in the horizontal plane at four different
heights. It can be seen that the microstructures in the horizontal plane and the vertical plane
display a similar tendency along the building direction: lathlike matrix structure, lamellar
structures, basketwave structures and colony alpha structures in the order of M1, M2, M3,
and M4, as shown in Figure 5-5 (c) to (f) respectively. In particular, the average width of
alpha lamellae at M1, M2 and M3 were about 1.89 μm, 1.80 μm and 0.89 μm respectively, as
shown in Fig.5-6, indicating that alpha grains from the areas at bottom to middle regions
have been refined by the thermal cycling. Due to accumulated heat serving as a preheating
CHAPTER 5 MATERIAL PROPERTIES
70
process, which contributes to the melting process of filler and base layer, the comparatively
higher temperature at M3 regions brings a higher density of dislocation nucleation, which can
additionally assist in the process of recrystallisation and creation of fine alpha grains.
(a) (b)
(c) (d)
CHAPTER 5 MATERIAL PROPERTIES
71
(e) (f)
Figure 5-5Optical micrographs of selected samples in different horizontal planes: (a)
schematic of the sample locations; (b) base metal; (c) M1; (d) M2; (e) M3; (f) M4.
Figure 5-6 Statistical distributions of the width of alpha lamellae for M1, M2 and M3
As a brief summary, the microstructural evolution of Ti6AL4V parts fabricated by
GT-WAAM can mainly be divided into four sections in order from bottom to top: lathlike
matrix structure, lamellar structure, basketwave structure and colony alpha structures.
CHAPTER 5 MATERIAL PROPERTIES
72
5.3.2 Phases characteristics and transformation
The same specimens used for micrographs were also used for XRD refinement
measurement to estimate the preferred orientation and the crystal phase of the deposited
Ti6Al4V alloy. As shown in Figure 5-7, the diffraction patterns for M1to M4 presents ten
broad peaks at 35.36°, 38.46°, 39.76°, 40.44°, 53.24°, 56.94°, 63.46°, 70.92°, 74.66°, 76.52°,
corresponding to α 101, α 002, β 110, α 101, β 101, β [125], α 112, α 101/β
211 and α 112 based on the analysis using Jade software.
Figure 5-7 XRD spectrums of M1, M2, M3 and M4
It is well known that the grains have their own preferred orientation of growth, namely
preferred-growth, dendritic-growth, and easy-growth directions. As can be seen from the Full
Width at Half Maximum (FWHM) of different peaks in the XRD patterns of Figure 5-7, for
the peaks of β 110 plane are almost invisible compared to α 101 peaks, which is
attributed to the fact that the β-Ti phase is usually too minor or absent to be detected by X-ray.
CHAPTER 5 MATERIAL PROPERTIES
73
With the increase of building height, the volume fractions of α 102 and α 110 decrease
significantly, indicating that more lamellar alpha phases are transformed into colony alpha
phase during the solidification process. It is worth mentioning that there is no obvious
preferred alignment direction for alpha lamellae in both vertical and horizontal planes due to
the generation of crystallographic texture along the columnar growth [74, 146].
5.3.3 Mechanical properties
At room temperature, the as-received base metal exhibits a yield strength (YS) of 884 ±
27 MPa, an ultimate tensile stress (UTS) of 995 ± 29 MPa and a total elongation of about
18.6%. Figure 5-8 shows the tensile test results of WAAM samples at eight locations along
the height of the wall. From locations S1 to S4, the YS varies slightly in the range of 892MPa
and 844 MPa, while the UTS varies between 958MPa and 1049MPa, which are comparable
to the base metal values. However, the mechanical properties are significantly degraded after
S4 (approximately 8mm above the substrate), where the interpass temperatures exceeds
200°C. The fracture of these samples occurred in the linear elastic stage, as shown in Figure
5-8b, which indicates an obvious decrease in the elongation and reduction of area in these
regions.
(a)
CHAPTER 5 MATERIAL PROPERTIES
74
(b)
(c)
(d)
Figure 5-8 Mechanical test results of the selected samples in horizontal plane of the
deposited wall: (a) the locations of tensile samples; (b) load-displacement relationships; (c)
ultimate tensile strength and yield strength; (d) elongation and reduction of area
CHAPTER 5 MATERIAL PROPERTIES
75
For Ti6Al4V alloy with a lamellar microstructure, the yield strength is primarily related
to the solid solution strengthening, including alpha plate boundary spacing and interstitial
solute level [147]. In the lower regions, the high interstitial content in both lathlike matrix
structure and lamellar structure is positive for improving YS and UTS. However, for the
regions where interpass temperature is above 200°C, more acicular and colony alpha were
formed, resulting in a decrease of interstitial content and alpha lamellae boundary spacing,
which deteriorates the mechanical strength. However, the precipitous drop in both YS and
UTS are mainly contributed by the oxidation due to the high interpass temperature. When the
interpass temperature is over 200°C, the trailing shield is not sufficient and a superficial
oxide coating can be clearly observed. As the interpass temperature increased, heavier
oxidation could be observed visually.
5.3.4 Facture behaviours
Figure 5-9 shows the appearance of tensile failure after uniaxial tension tests. It can be
recognised that the samples from S1 to S4 present obvious lamellar tearing and neck
contraction in fractural locations, while the samples from S5 to S8 break with no deformation,
which is in line with the elongation results in the previous section.
Figure 5-10 further displays the fractographs of these tensile specimens. The shear lips
can be clearly observed on samples S1 to S4, especially for S1 with a large number of
dimples in lower regions, which illustrates a typical ductile fracture. Due to rapid heat
dissipation into the substrate in the first few layers, alpha lamellae structures were formed.
The areas of interphases decreased, resulting in the promotion of void growth [148] and
generation of a number of dimples. However, samples S5 to S8 show more brittle cleavage
CHAPTER 5 MATERIAL PROPERTIES
76
morphology in response to fracture, which is contributed from the increase of colony alpha
structures decorated within beta grains and pinning effects with carbon and oxygen solute.
According to a study of the fracture characteristics for Ti6Al4V parts fabricated using LSM,
it was found that the tensile fracture mainly occurred following the beta grain boundary as
well as concomitantly generated internal cracks [149]. As shown in Figure 5-11, the cracks
that appear in S7 indicate that intercrystalline fracture occurs and the tensile strength is
greatly reduced. Furthermore, the alpha phases decorated at the prior beta grain boundary
also have significant effects on the fracture features [147]. To be more specific, the alpha
grain boundary is much softer than the beta grain, which results in the interface at phase
boundaries being more brittle in the fracture tests[149]. Therefore, the regions where more
beta grains have formed exhibit cleavage fracture features, and this also explains why the
elongation of these regions is very low with fracture failure.
In addition, through EDS analysis, two important interstitial elements, carbon and
oxygen, are observed in the fractural surface of upper regions, which indicates potential
chemical reactions of Ti with O2 and CO2 from the surroundings. This induces the pinning
effects with carbon and oxygen solute atoms in the microstructures [150], leading to
unacceptably low ductility in the influenced regions as well as fracture failure at the linear
elastic stage.
The critical ductility factor, reduction of area, can also be used to distinguish the fracture
features. For a tensile specimen with reduction of area greater than 5%, it would be defined as
ductile fracture; otherwise it would be classified as brittle fracture. From Figure 9d, it can be
seen that samples S1 to S4 undergo more than 5% reduction of area, which indicates that
CHAPTER 5 MATERIAL PROPERTIES
77
these locations within the fabricated wall should exhibit ductile fractures, as compared to
samples S5 to S8. These results also explain why regions along the building direction have
different fracture features. It is evident that during the manufacture of Ti6Al4V with the
GT-WAAM process using localized gas shielding, the changes in oxidation and
microstructural evolution that are caused by altered thermal behaviours have a significant
negative impact on the mechanical properties of fabricated parts.
Figure 5-9The fracture appearance of the tensile specimens
(a) (b)
CHAPTER 5 MATERIAL PROPERTIES
78
(c) (d)
(e) (f)
(g) (h)
Figure 5-10 High-magnification fractographs of corresponding tensile samples: (a) S1; (b)
S2; (c) S3; (d) S4; (e) S5; (f) S6; (g) S7; (h) S8.
CHAPTER 5 MATERIAL PROPERTIES
79
Figure 5-11 Axial crack in fracture surface of S7
5.4 Conclusion
In this section, Ti6AL4V alloy was fabricated by the GT-WAAM process using only
localized gas shielding. Based on the in-situ temperature measurement at each layer, the heat
accumulation during the deposition process was calculated. OM, XRD, SEM, EDS and
standard mechanical test investigations were performed, and demonstrated that the
microstructural morphology, crystalline phase, mechanical properties and fracture feature
features varies are all changed as heat accumulates along the building direction. The findings
include:
(1) To obtain favorable mechanical properties, limiting interpass temperature to 200°C is
desirable for Ti6Al4V fabrication when using WAAM technology with localized gas
shielding.
(2) The microstructural evolution influences the mechanical properties and fracture
features of the as-fabricated part through the phase transition in the development of lathlike
matrix structure, lamellar structure, basketwave structure and colony alpha structure along the
CHAPTER 5 MATERIAL PROPERTIES
80
building height. The control of colony alpha structure formation is beneficial to the
improvement of ductility.
(3) The discrepancy of the mechanical properties in different locations along the
deposited wall can be explained by the dominant effects of oxidation and the partial influence
from microstructural evolution that are collectively determined by the layer’s interpass
temperature.
In summary, a better understanding of the effects of heat accumulation on material
properties during the WAAM process will benefit future process control, improvement, and
optimization.
CHAPTER 6 CORROSION RESISTANCE
81
Chapter 6 CORROSION RESISTANCE
6.1 Introduction
For many applications, corrosion resistance is a primary consideration because it often
determines the service life of the component. A number of studies have assessed the corrosion
behaviour of Ti-6Al-4V under various conditions. Dai et al. [151] evaluated the corrosion
resistance of selective laser melted (SLM) Ti-6Al-4V in NaCl solution, and found that the
SLM-produced samples had poorer resistance to corrosion than commercial grade 5 wrought
alloy due to the acicular α’ martensite formed within the microstructures. The authors further
reported that anisotropic corrosion behaviour exists in SLM-fabricated Ti-6Al-4V alloy[152].
In 1 M HCI solution, corrosion resistance in the planes perpendicular to the build direction is
superior to that in planes parallel to the build direction, while in 3.5 wt.% NaCl solution, the
corrosion resistances show only a very slight difference. Recently, a group of electrochemical
corrosion tests in NaCl solution were carried out by Yang et al.[153] to compare the corrosion
resistance of Ti-6Al-4V specimens that were processed by SLM, SLM followed by heat
treatment (SLM-HT), WAAM and also traditional rolling conditions. The SLM-HT sample
exhibited the highest corrosion resistance, followed by rolled, WAAM and finally SLM. The
differences in performance were attributed to the formation of distinctive microstructures by
the different manufacturing processes. In the cause of WAAM-produced Ti-6Al-4V, the
corrosion mechanism is still not well understood due to its complex microstructural
distribution and limited information reported in earlier literature. It is believed that the
electrochemical corrosion resistance of WAAM-fabricated Ti-6Al-4V has anisotropic
CHAPTER 6 CORROSION RESISTANCE
82
characteristics due to its directional microstructures. A better understanding of the underlying
mechanism may be beneficial from the perspective of optimizing microstructural control in
the WAAM process.
In this section, a comprehensive investigation on the electrochemical corrosion
behaviour of WAAM-fabricated Ti-6Al-4V part has been conducted by means of optical
microscopy (OM), X-ray diffraction (XRD), hardness testing, potentiodynamic polarization
and electrochemical impedance spectroscopy (EIS) analysis. Specimens were produced from
the fabricated part in orientations parallel and perpendicular to the build direction, to assess
the influence of directional microstructure on corrosion behaviour. Samples were also taken
from the build substrate, to assess the comparative corrosion behaviour of conventional
roll-processed or wrought Ti-6Al-4V. The findings provide an insight into the corrosion
mechanism of WAAM fabricated Ti-6Al-4V, which gives direction to future improvements of
process control and optimization.
6.2 Experimental Procedures
6.2.1 Sample and solution preparation
The apparatus used in this study mainly consists of a 200A-rated GTAW power source,
“cold” wire feeder, water cooling unit and travel mechanism, as shown in Figure 6-1. The
feedstock was commercial ASTM B863 grade 5 Ti-6Al-4V alloy wire with a diameter of
1.2mm deposited onto a Ti-6Al-4V substrate with dimensions of 200 mm × 150 mm × 6 mm
to ASTM B265 specification. A straight wall structure of 150 mm in length, approximately 10
mm in width and fifteen layers in height was fabricated with the prepared GT-WAAM system
CHAPTER 6 CORROSION RESISTANCE
83
using localized gas shielding. The process parameters are provided in Table 6-1. Welding
grade argon (99.995% purity) was used for both GTAW torch and trailing shield. After
deposition, the fabricated wall was sliced into two groups of specimens: vertical planes (VP)
and horizontal planes (HP), respectively, as shown in Figure 6-2. In addition, the same size
of specimen was cut from the substrate base metal (BM) and used as a third investigation
group.
Figure 6-1 Schematic diagram of GT-WAAM system
Table 6-1 Process parameters for GT-WAAM deposition
Process parameters Details
Deposition current 110 A
Arc voltage 12 V
Travel speed 95 mm/min
Wire feed speed 1000 mm/min
Distance between the electrode and workpiece 3 mm
Angle between the electrode and the filler wire 60°
Flow rate of argon in GTAW torch 10 L/ min
Flow rate of argon in trailing shield 10 L/ min
Post flow duration 35 seconds
Dwell time between layers 125 seconds
CHAPTER 6 CORROSION RESISTANCE
84
Figure 6-2 Three-dimensional diagram of WAAM-fabricated Ti-6Al-4V wall showing
orientation of specimen planes
6.2.2 Material analysis
The extracted specimens for microstructural observation were hot mounted, ground and
polished according to standard procedures, and then etched in Kroll’s reagent containing 2%
vol HF, 6 vol% HNO3 with balance H2O. The microstructures were observed using a Leica
DMR optical microscope (OM). Qualitative micro-analysis of phases constituents were
determined by grazing-incidence X-ray diffraction (XRD) with a GBC MMA X-ray
diffractometer using monochromatic Cu Kα radiation at an accelerating voltage of 35 kV and
a current of 28.6 mA in scanning steps of 0.02º and a scanning speed of 4 º/min in the range
of 30 ~ 90º. Vickers hardness testing was performed with a DuraScan 70 automatic hardness
tester using a test load of 0.98 N and a dwell time of 15 s according to ASTM: F2924. The
average value of ten measurements was taken as the hardness for each specimen.
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6.2.3 Electrochemical measurements
All the samples to corrosion test were subjected to grinding with 220 # SiC paper for 5
minutes, followed by polishing with different size of clothes from 9 μm down to 1μm
(standard metallographic method for Ti6Al4V alloys) using Struers Tegrapol 21 automatic
grinder and polisher. The electrochemical tests included open circuit potential (OCP),
potentiodynamic polarization and electrochemical impedance spectroscopy (EIS). All the
electrochemical experiments were conducted on a CH Instruments 600C electrochemical
station with a conventional three-electrode cell in 3.5 wt.% NaCl solution at room
temperature. A platinum sheet and saturated calomel electrode (SCE) were used as counter
electrode and reference, respectively. The electrochemical active surface area of working
electrodes was kept at 0.2 cm2. Before polarization and EIS tests, the specimens were
immersed in solution for sufficient time to obtain a stable OCP. EIS was performed in the
frequency range of 0.01Hz to 100 kHz with an amplitude of 10 mV. Subsequently, the
potentiodynamic polarization experiments were carried out at ±150 mV versus OCP at a
sweep rate of 0.5 mV/s. Strong polarization tests were carried out in the potential range of −1
V to + 2 V at a scan rate of 5 mV/s in 3.5 wt.% NaCl solution. The electrochemical
measurements were divided into three groups: vertical plane (VP), horizontal plane (HP) and
base metal (BM), each group having three specimens selected from identical planes for
checking reproducibility of data.
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6.3 Results
6.3.1 Microstructural studies
6.3.1.1 Metallographic microstructure
Figure 6-3 shows the optical micrographs of BM, VP and HP regions of the fabricated
specimen. The main microstructure of wrought Ti-6Al-4V alloy presents a mixed α/β
equiaxed phase structure distributed uniformly. For the WAAM-fabricated Ti-6Al-4V
samples, a secondary α morphology (acicular martensite α’) is preferentially formed in HP
regions, while a fully lamellar α morphology forms in VP regions, interwoven with a
Widmanstätten structure inside the prior-grain boundary α. The microstructure of
WAAM-fabricated Ti-6Al-4V is inhomogeneous in both HP and VP regions, which is likely
to produce non-uniform properties and performance.
(a)
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(b)
(c)
Figure 6-3 Optical micrographs of three selected specimens: (a) BM region; (b) HP region; (c)
VP region.
6.3.1.2 Phase composition
The same specimens used for micrographs were also used for XRD measurement to
determine the preferred orientation and the crystal phase of the alloy. As shown in Figure 6-4,
the diffraction pattern for the BM, HP and VP samples presents eight broad peaks at 35.36°,
38.46°, 40.44°, 53.24°, 63.46°, 70.92°, 76.52°, 87.42° based on analysis using Jade software.
CHAPTER 6 CORROSION RESISTANCE
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Figure 6-4 The XRD spectrums of three selected specimen from different planes
As can be seen from Figure 6-4, the microstructure of the BM sample presents a large
number of α-Ti phases. These are replaced by predominantly α’-Ti phases in the HP of
WAAM-fabricated Ti-6Al4V samples. It should be pointed out that only small amount of β-Ti
phases are observed in the samples, which is attributed to the fact that the β-Ti phase is
usually too minor or absent to be detected by X-ray [154]. Although both vertical plane and
horizontal plane samples of WAAM-fabricated Ti-6Al-4V have similar microstructural
composition, the relative fractions of the phases are still different, especially for the
martensitic α’ 102 and 110 crystallographic plane, displaying a large volume fraction in
HP regions. This indicates that α’ 102 and 110 planes are the preferred orientations of
growth in the HP regions, which ultimately causes the microstructural evolution to vary along
different directions.
6.3.1.3 Mechanical properties
The average hardness distribution of the specimens in the three selected regions for the
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corrosion test is shown in Figure 6-5. As can be seen, both VP and HP regions have lower
average hardness compared to BM regions, possibly due to the microstructural changes in
these regions. Even with such a similar microstructural composition between HP and VP
regions, discernible differences in average hardness still exist, displaying higher values in VP
regions than HP regions. This is attributed mainly to the influences of grain size and phase
orientation on dislocation distributions in the different planes.
Figure 6-5 Average hardness values of three groups of specimens
6.3.2. Electrochemical studies
6.3.2.1 Open circuit potential (OCP)
Open circuit potential (OCP) is considered as an important parameter in estimating the
tendency of metal corrosion. As shown in Figure 6-6, before performing other
electrochemical tests, the variation of EOCP with time was measured by immersing the
samples in 3.5% NaCl solution. The EOCP values of the as-received specimens increase with
time, which indicates that the electrochemical process is controlled by the anodic reaction and
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a protective oxide film starts to grow on the electrode surface. During immersion in the
corrosive media, the newly formed oxide film may be dissolved, and then formed locally
again. This simultaneous competition between film formation and dissolution brings an
unstable electrochemical process to the electrode surface, consequently leading to fluctuating
EOCP values. In comparison to BM and VP samples, HP samples exhibit the most variation of
EOCP values owing to less homogeneous microstructures. After a certain immersion time, the
oxide film on the electrode surface begins to stabilize and therefore, the EOCP values approach
a constant. As the immersion time approaches 60000 s, EOCP values of all investigated
samples stabilized at around −0.1 V, with slight differences between fabricated samples and
wrought metal. It is well recognized that the nobler, or more positive, OCP indicates greater
stability. Hence from Figure 6-6, the studied samples exhibit a tendency over time toward
decreasing corrosion tendency.
Figure 6-6 Open circuit potential (vs. SCE) of as-received samples in 3.5% NaCl solution.
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6.3.2.2 Potentiodynamic polarization
In this study, the potentiodynamic polarization is used to calculate the kinetic parameters,
including corrosion current density (Icorr), corrosion potential (Ecorr) and Tafel slopes (ba and
bc). Tafel extrapolation method was adopted as one of the most commonly used method for
titanium corrosion test, including Ti6Al4V [155] and other kinds of titanium-based alloy
[156]. Recently, there was report on investigation of corrosion resistance by the method of
Tafel extrapolation on Ti6Al4V alloy fabricated by electron beam melting [157]. Figure 6-7
shows the Tafel extrapolation results of the investigated specimens in 3.5 wt.% NaCl solution,
and Table 6-2 further details the analysis results. As can be seen, HP and VP samples display
lower average Ecorr values (−0.138 V and −0.131 V respectively) compared with the BM
sample (−0.117 V), indicating slightly less corrosion resistance for both of them. Regarding
the corrosion current density (Icorr) which generally reflects the corrosion rate for the
materials, HP samples display the highest average Icorr value (16.0 nA/cm2), followed by VP
samples (14.2 nA/cm2) and BM samples (12.1 nA/cm2). Larger Icorr values indicate higher
corrosion rates, so the BM samples show a better corrosion resistance compared to the
WAAM-fabricated Ti-6Al-4V samples in NaCl solution. By this measure, the VP samples
exhibit a lower corrosion rate than HP samples, in general agreement with Ecorr values.
Moreover, the BM samples demonstrate better repeatability in both Ecorr and Icorr values due
to more uniform microstructure, while VP and HP samples show greater variation, which is
not only related to microstructural heterogeneity but also influenced by grain size and phase
orientation. In terms of linear polarization resistance (Rp'), both sets of WAAM-fabricated
Ti-6Al-4V samples have lower average values (2.27 M cm2 for VP, 2.23 M cm2 for HP)
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in comparison to the wrought substrate (2.49 M cm2 for BM), suggesting poorer corrosion
resistance, which is consistent with the results for Ecorr and Icorr.
Figure 6-7 Potentiodynamic polarization plots of specimens
Table 6-2 Electrochemical kinetics parameters for test samples
Specimen Ecorr
(V/SCE)
Icorr
(nA/cm2)
ba
(1/V)
-bc
(1/V)
Rp'
(M cm2)
V1 -0.224 11.6 5.19 10.11 2.44
V2 -0.018 20.8 4.83 7.47 1.70
V3 -0.153 10.3 3.77 12.12 2.66
H1 -0.137 19.5 4.32 4.44 2.54
H2 -0.15 20.0 4.86 11.89 1.28
H3 -0.108 8.33 5.99 12.22 2.86
B1 -0.129 15.5 4.06 7.97 2.32
B2 -0.106 8.34 6.78 12.88 2.66
B3 -0.117 12.6 4.61 8.62 2.50
Figure 6-8 shows the strong polarization curves of WAAM-fabricated Ti-6Al-4V and
wrought Ti-6Al-4V. These tests were carried out in the potential range of −1 V to + 2V at a
scan rate of 5 mV/s in 3.5 wt.% NaCl solution. According to Fig.8, all the investigated
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93
specimens are characterized by one obvious passivation region where the increase of current
density with potential is inhibited. When the corrosion potential (Ecorr) of the electrode
surface exceeds 0 V, the potentiodynamic curves deviate from the Tafel anodic line, meaning
passivation behaviour begins to occur. In the passive regions, there is a slight difference in
passive potential range (ΔE) between WAAM fabricated Ti-6Al-4V and wrought Ti-6Al-4V.
To further evaluate the passivation process, the passive current density (ip) can be identified
in the magnified drawing in Figure 6-8. If the electrode surface shows higher passive current
density, it is less likely to be passivated. As can be seen, HP samples exhibit the highest ip
value, followed by that of VP and BM samples, which suggests the HP samples exhibit a
comparatively larger corrosion rate even in passive regions.
Figure 6-8 Strong polarization curves for WAAM-fabricated Ti-6Al-4V and ASTM standard
Grade 5 alloy in 3.5wt% NaCl solution. The inset shows passive region from 0 V to 2 V.
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6.3.2.3 Electrochemical impedance spectroscopy (EIS)
Figure 6-9 shows the Nyquist and Bode plots of the VP, HP and BM specimens in 3.5wt%
NaCl solution. The equivalent circuit for fitting the EIS data is shown as an inset in Figure
6-9 (a). For all the three sets of samples, since the Nyquist diagram (Figure 6-9 (a)) exhibits
only an obvious capacitive loop, a Randles circuit with one time constant was selected as the
equivalent circuit to fit EIS results. In the equivalent circuit, Rs is the solution resistance, Rf
refers to film resistance of the passive film, and the constant phase element (CPE) represents
film capacitance, which can be expressed as [158]:
ZCPE=Y0−1(jω) −n (6-1)
where Y0 is true capacitance of the passivation layer, j is imaginary unit, ω is angular
frequency and n is related to the constant phase angle[159, 160]. The detailed EIS fitting data
are summarized in Table 6-3. A smaller value of χ2 indicates a better fit between the
equivalent circuit and experimental results.
(a)
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(b)
Figure 6-9 EIS results of test samples: (a) Nyquist plots with inset showing the equivalent
circuit: (b) Bode plots.
According to the data listed in Table 6-3, for HP, VP and BM samples, the solution
resistance (Rs) values in 3.5. wt.% NaCl solution are very similar, however there are
significant differences in passive film resistance (Rf).While BM samples display the highest
value of Rf (average value 1.68 MΩ cm2), VP samples exhibit a significantly lower average
Rf value (195 kΩ cm2), but distinctively higher than that of HP samples (average 174 kΩ
cm2). Generally, a higher Rf value of passive film implies better corrosion resistance. Based
on these results, WAAM-fabricated Ti-6Al-4V has an inferior corrosion resistance when
compared to wrought Ti-6Al-4V. The value of n represents the degree to which the constant
phase element (CPE) acts as a purely capacitive element, with diminished contribution of a
CHAPTER 6 CORROSION RESISTANCE
96
resistive component. All the investigated Ti-6Al-4V specimens have n values close to the
number 1.0, so the CPE in the equivalent circuit closely resembles an ideal capacitor. In order
to further evaluate the metal surface condition immersed in NaCl solution, the passive film
thickness can be estimated using the parallel plate capacitor equation:
0 Ad
C
= (6-2)
where d presents the thickness of the passive layer, ε is the relative permittivity of the film, ε0
is permittivity of vacuum, A is the exposed area of the working electrode and C is related to
the value of Y0, CPE. According to this relationship, the passive film produced on the surface
of BM specimens (average Yo 25.9) is expected to be thicker on average than that of VP
specimens (average Yo 29.5) HP specimens (average Yo 31.8) when evaluated in 3.5 wt.%
NaCl solution.
Table 6-3 EIS parameters of equivalent circuit for test samples
Specimen Rs
( cm2)
CPE, Y0
(10-6S Secn cm-2)
n
(1/V)
Rf
(M cm2)
χ2
(10-3)
d
(nm)
V1 30.4 28.5 0.895 0.137 2.93 61.1
V2 31.0 37.0 0.847 0.148 4.73 47.1
V3 32.0 22.9 0.899 0.301 5.16 75.9
H1 30.2 24.4 0.913 0.173 2.00 71.3
H2 30.8 48.7 0.832 0.125 2.30 35.8
H3 29.6 22.3 0.916 0.185 1.93 77.9
B1 29.5 19.2 0.913 1.47 2.51 90.8
B2 31.2 27.1 0.889 2.04 2.53 64.2
B3 30.5 24.3 0.908 1.52 2.83 71.6
6.4 Discussion
6.4.1. Microstructural evolution
For the VP, HP and BM regions, the main variation in microstructure is the generation of
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97
martensite lamellae α with Widmanstätten structure and acicular α’ structure due to the
differences in thermal behaviour at each location. As layers are successively deposited during
the WAAM process, the cooling rate progressively decreases for higher layers, and this
contributes to the variation in microstructural evolution.
The conductive thermal resistance from the uppermost deposited layer to the substrate is
significantly increased with increasing wall height. As titanium alloys exhibit poor thermal
conductivity, a decreasing fraction of the heat input from the gas tungsten arc welding
process is conducted away through the substrate, while an increasing amount of heat is
dissipated to the surrounding atmosphere via convection and radiation, which leads to a
slower cooling rate in the building direction (VP) than in the travel direction (HP) [141].
Therefore, as shown in Figure 6-10, upon high cooling rates from the β phase field, the
transformation from β phase to α phase occurs by a displacive transformation to form acicular
martensite structures along the travel direction (HP). Generally, these acicular martensite
structures consist of mixtures of individual α’ plates each belonging to different Burger
orientations [161] (Figure 6-10b), and will lose their hexagonal structure to a orthorhombic
α’’ structure with the β stabilizing solute increase [142]. Conversely, due to slow cooling rates
along the building direction (VP), the α phase nucleates preferentially at the prior β grain
boundaries at first, forming a continuous layer named the grain boundary α (αGB), which
always exhibits a Burgers orientation related with one of the prior β grains [162]. Afterwards,
the nucleation of α lamellae occurs in the grain interior so that basketwave Widmanstätten α
forms in pseudo-random fashion within the remaining β structure (Figure 6-10c).
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98
Figure 6-10 Schematic of microstructural evolution for WAAM-fabricated of Ti-6Al-4V: (a)
Continuous cooling diagram for Ti–6Al–4V β-solution treated at 1050°C for 30 min (Ms
temperature due to Majdic and Ziegler) [142], (b) HP regions, (c) VP regions.
Referring to the hardness measurements in Figure 6-5, higher values are measured in
the BM samples, followed by VP and then HP. These disparities are obviously associated to
the formation of α and α´ martensite microstructure. Generally, the martensitic lamellae α at
VP show a higher hardness value in comparison with martensite α´ at HP due to its
comparatively larger grain size and the formation of grain boundaries. For Ti-6Al-4V α + β
structures, the Vickers hardness decreases with grain size and the thickness of lamellar α
phase increasing due to its positively proportional to strength of alloy[163]. Therefore,
slightly high hardness value at VP is also beneficial to corrosion resistance improvement.
6.4.2 Electrochemical evaluation
The findings of the electrochemical tests in 3.5 wt.% NaCl solution are useful in
assessing the comparative corrosion resistance of the investigated specimens. As described in
Figure 6-6, rising OCP values indicate that a stable passive film forms on the electrode
surface and these passive films play a key role in protecting inner metal from further
CHAPTER 6 CORROSION RESISTANCE
99
corrosion. Potentiodynamic polarization and EIS results suggest that the wrought Ti-6Al-4V
possesses higher corrosion resistance than WAAM-fabricated Ti-6Al-4V in which inner
planes parallel to the build direction (VP) show slightly better corrosion properties on
average than those planes that are perpendicular to the build direction (HP). The
electrochemical test results that are critical for an in-depth understanding of the corrosion
resistance of WAAM-fabricated Ti-6Al-4V and wrought Ti-6Al-4V are shown in Figure
6-11.
In general, material properties are closely related to microstructure, and this is
particularly the case for corrosion resistance. For WAAM-fabricated Ti-6Al-4V, the formed
lamellae α and acicular α’ grains are always present as a non-equilibrium structure that is
believed to have a high energy state due to inadequate atomic diffusion during solidification.
This attribute ultimately leads to unstable microstructures[74] and produces inferior corrosion
resistance, as has been measured in this study using in 3.5 wt. % NaCl solution. Also, for the
selective corrosion of α phase and α’ phase, the dissolution rate of α phase is lower than that
of α’ phase on the surface of specimens at open circuit potential due to the comparatively
hard grain [164]. Therefore, the HP regions with acicular α’ phase, which is in a “higher
energy state”and metastable, has inferior corrosion resistance compared to the VP regions
with lamellar α structure. It needs to be noted that the greatest difference in these two types of
microstructures is their grain size. In corrosive media, both of these massive grains perform
as effective buffer films, contributing to the reduction of the galvanic effect between α and β
phase. However, in comparison with acicular α’ structure, the lamellar α structure possesses a
higher density of nucleation sites for passivation during the corrosion process[165], which
CHAPTER 6 CORROSION RESISTANCE
100
leads to a thicker passive film and lower corrosion rate[166]. Accordingly, an averagely better
corrosion resistance is obtained in VP regions of WAAM-fabricated Ti-6Al-4V.
For Ti-6Al-4V alloy with α + β microstructure, the Al and V elements are usually acting
as respective stabilizers for α phase and β phase, and these segregation of elements always
leads to an intrinsic difference in the corrosion resistance of α phase and β phases [167]. In
corrosive solutions, if more preferential dissolution of α phases exist, the corrosion rate will
be accelerated [168], while β phase is comparatively resistant to corrosion and improves
corrosion resistance. Based on the XRD results, although only small amounts of β-Ti phases
were detected in all the investigated specimens, difference in the β-Ti phase content still exist.
It is observed that more β grains are detected in BM regions, which means enhanced
corrosion resistance. Meanwhile, this is also one of the reasons that why the BM samples
have higher hardness values than the WAAM-fabricated areas.
In addition, it is worth noting that the microstructural uniformity of investigated
specimens not only has a direct effect on the corrosion resistance of the metal, but also brings
a serious impact to the repeatability of electrochemical results. In this study, the BM samples
have more uniform microstructures than WAAM-fabricated samples, and that is reflected in
the consistency and repeatability of measurements. However, even in adjacent regions of
fabricated Ti-6Al-4V, the examined microstructure is markedly inhomogeneous, producing
significant fluctuations in the electrochemical test results. The surface of samples subtracted
from vertical direction (VP), like geological layers, consists of a greater variation of
microstructures than that of HP samples. As the active area of corrosion electrodes is
relatively small, the greater heterogeneity of VP samples results in the greater variation of
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101
electrochemical corrosion test results. Regardless of the individual, localized influences
created by microstructural non-uniformities, the overall findings remain relevant. The
anisotropic microstructure resulting from the WAAM processing of Ti-6Al-4V has a
significant effect on the corrosion performance. Any future improvement to the process that
aims to improve this aspect must address the mechanisms that produce microstructural
anisotropy. It is well acceptable that the microstructural evolution is essentially dependent on
the thermal history of the process that involves alternate re-heating and re-cooling cycles.
That is to say, if the thermal state during deposition is properly controlled, the desired
microstructures and more uniform properties with acceptable performance can be achieved.
Based on this interpretation, an innovative wire arc additive manufacturing (WAAM) process
with interpass heat sink will be further developed, trying to obtain more uniform
microstructures and properties through the reasonable control on thermal state to fabricated
part.
Figure 6-11 The comparison in corrosion resistance of WAAM-fabricated Ti-6Al-4V and
wrought base metal in 3.5% NaCl solution.
CHAPTER 6 CORROSION RESISTANCE
102
6.5 Conclusion
The effects of directional microstructures on the corrosion properties of Ti-6Al-4V parts
fabricated with wire are additive manufacturing (WAAM) technology have been investigated,
and these corrosion properties have been compared against the performance of ASTM
standard wrought-Ti-6Al-4 alloy. As the WAAM component is built up from the substrate, the
conductive thermal resistance through the deposited material varies along different
orientations. A relatively slower cooling rate occurs in the build direction (VP-vertical plane),
contributing to the formation of large and hard lamellar α grains that are interwoven with a
Widmanstätten structure. Meanwhile, a comparatively faster cooling rate occurs in the travel
direction (HP-horizontal plane), producing smaller and softer α’ grains. The resulting
differences in phase structure, grain size and orientation creates a greater susceptibility to
corrosion in the HP orientation compared to the VP orientation, as found by testing in 3.5%
NaCl solution. Furthermore, WAAM-fabricated Ti-6Al-4V in both VP and HP orientations
exhibits a lower passive film thickness than wrought Ti-6Al-4V, indicating higher corrosion
rates and inferior corrosion resistance. In addition, the non-uniform microstructures produced
in WAAM fabrication are not only responsible for the anisotropic corrosion behaviour, but
also lead to greater variation in corrosion test results. To address this shortfall in the potential
service life of additively manufactured Ti-6Al-4V components that are to be used in corrosive
environments, further improvement of the WAAM process are needed, to address the
mechanisms that produce microstructural anisotropy.
CHAPTER 7 THE FORCED INTERPASS COOLING
103
Chapter 7 THE FORCED INTERPASS COOLING
7.1 Introduction
For a Ti6Al4V component that is produced by WAAM, the microstructures are complex,
often varying spatially within the deposition due to its complex thermal history that involves
alternate re-heating and re-cooling cycles [169]. Even for adjacent regions within a deposit,
differences in microstructure still exist, ultimately bringing inhomogeneous material
performance to the component. To date, numerous investigations have been carried out to
reveal microstructural formation mechanisms through the investigation of thermal history in
WAAM deposition, with the aim of offering some helpful directions to solve these problems.
For example, Wu et al. [141] observed that heat accumulation during additive fabrication has
significant effects on the stability of gas tungsten wire arc additive manufactured Ti6Al4V
alloy, which demonstrates the importance of in-process temperature monitoring and control.
Suryakumar, et al. [170] claimed that depending on the thermal cycles in WAAM-fabricated
steel, coarse grained microstructures with inhomogeneous hardness are formed along the
building direction. These findings suggest that there is a tremendous need for microstructural
optimization through improvements in process control of additive manufacturing.
As mentioned above, the solidification microstructure, including grain size and
morphology, is essentially dependent on the thermal history during the manufacturing process.
That is to say, if the thermal state during deposition is properly controlled, the desired
microstructures and resulting mechanical properties with acceptable performance can be
achieved. Based on this interpretation, recent studies by Henckell et al. [171] have attempted
CHAPTER 7 THE FORCED INTERPASS COOLING
104
to employ additional cooling gas for WAAM fabrication of low alloy steel components.
Interestingly, applying an additional post-weld cooling gas of nitrogen with 5% hydrogen
improves both the layer geometry and mechanical properties through grain refinement and
homogenous hardness. Additionally, the different gas types used in deposition produce
varying material properties, and further investigations were proposed that would employ
carbon dioxide or pure nitrogen. It is well known that CO2 gas cooling is commonly
employed in Low Stress No Distortion (LSND) welding to reduce residual stress and
deformation. According to Holder et al. [172], the heat affect zones of DH-36 steel welds
show significant changes in hardness when subjected to conventional gas metal arc welding
and LSND welding processes. Based on such findings, if CO2 gas is used to dynamically cool
the deposits during the WAAM process, it is possible that similar effects on material
properties can be obtained.
In this study, gas tungsten arc welding based WAAM (GT-WAAM) with forced interpass
cooling using compressed CO2 gas is applied to Ti6Al4V, with the aim of improving material
properties. The effects on surface oxidation, geometrical features, microstructural evolution,
and mechanical properties are investigated through the use of WAAM with and without CO2
gas cooling. Firstly, the interpass temperature is controlled (i.e. allowed to decrease through
natural convection and conduction before deposition of the next layer) to a number of
constant values, aiming at an understanding of the effects of interpass temperature on the
targeted deposition properties. Subsequently, forced interpass cooling is implemented, aimed
at an improvement of material properties. A detailed analysis and discussion of the results is
used to assess the possibilities for controlling the properties of WAAM-fabricated Ti6Al4V
CHAPTER 7 THE FORCED INTERPASS COOLING
105
by using this method.
7.2 Experimental Procedures
7.2.1 Experiment setup
The apparatus consists of a 200A-rated GTAW power source, “cold” wire feeder, water
cooling unit and travel mechanism, as shown in Figure 7-1. An IR pyrometer was used to
measure the in-situ interpass temperature, thereby controlling the dwell time between
deposition of layers. The feedstock used in this study was ASTM B863 grade 5 Ti6Al4V wire
with a diameter of 1.2mm, deposited onto a Ti6Al4V substrate with dimensions of 200 mm ×
150 mm × 6 mm that conforms to ASTM B265 specification. Three straight wall structures
were fabricated at respective interpass temperatures of 100 °C, 200 °C, and 300 °C using the
GT-WAAM system with localized gas shielding. The process parameters are provided in
Table 7-1. Welding grade argon (99.995% purity) was used as the shielding gas for both
GTAW torch and trailing shield.
Figure 7-1 Schematic diagram of the GT-WAAM system
CHAPTER 7 THE FORCED INTERPASS COOLING
106
Table 7-1 Process parameters for WAAM deposition
Process parameters Details
Deposition power Current 110A
Arc voltage 13V
Speed Travel speed 95mm/min
Wire feed speed 820mm/min
Distance and angle Electrode to workpiece 3 mm
Electrode to filler wire 60°
Flow rate(argon) GTAW torch 15 L/ min
Trailing shield 10 L/ min
Time Post flow duration 35 seconds
7.2.2 Forced interpass cooling
In order to evaluate the effectiveness of forced interpass cooling, a
commercially-available CO2 cooling spray nozzle was fitted to the GTAW torch as shown in
Figure 7-2. A flow rate of 215g/min was used to deliver a stable CO2 cooling stream,
according to the manufacturer’s specifications. The cooling spray was initiated after
deposition of each layer was completed in order to avoid arc disruption caused by the
relatively turbulent CO2 flow. Cooling flow was stopped when the interpass temperature
decreased to room temperature.
Figure 7-2 Schematic diagram of the GT-WAAM deposition system with forced
interpass cooling
CHAPTER 7 THE FORCED INTERPASS COOLING
107
7.2.3 Material characterization techniques
Both metallographic specimens and mechanical test samples were extracted from the
fabricated Ti6Al4V walls for further investigation. The metallographic specimens were hot
mounted, ground and polished according to standard procedures, and then etched in Kroll’s
reagent containing 2% vol HF, 6% vol HNO3 with balance H2O. The macrostructures were
examined using a Leica M205A deep field stereoscopic microscope and the microstructures
were observed using a Leica DMR optical microscope (OM). Vickers hardness testing was
performed with a DuraScan 70 automatic hardness tester using a test load of 100g and a dwell
time of 15s according to ASTM: F2924 standard. Tensile tests were carried out at room
temperature at a constant crosshead displacement rate of 0.4mm/min using a MTS370
universal testing machine. The tensile test samples extracted from the fabricated walls are
shown in Figure 3.
Figure 7-3 Schematic diagram of extracted tensile samples
7.3 Results
7.3.1 Deposition geometry
Figure 7-4 shows the dimensions of WAAM-fabricated Ti6Al4V walls at different
CHAPTER 7 THE FORCED INTERPASS COOLING
108
interpass temperature. When the interpass temperature increases from 100°C to 300°C, the
height of the deposited wall is reduced by 11.7%, whereas its width increases to 12.2%.
Interpass temperature has a significant effect on build geometry, even though all other
process parameters are unchanged, including the number of layers. Also, the dimensions of
the wall deposited using forced interpass cooling are almost identical to those of the wall
built with an interpass temperature of 100 °C. This indicates that forced interpass cooling not
only reduces the build time, but also has the advantage of maintaining deposition accuracy.
Figure 7-4 Effect of interpass temperature on wall dimensions
7.3.2 Surface oxidation
The images in Figure 7-5 show that different degrees of surface oxidation can be
observed at various interpass temperatures. For Figure 7-5 (a), the interpass temperature is
allowed to return to 100 °C through natural convective and conductive cooling. The initially
low temperature at the start of the next deposited layer essentially eliminates the effects of
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heat input from the newly deposited layer, so the surface of fabrication shows a clean surface
with only a very slight sign of oxidation as indicated by the light straw color. As the interpass
temperature is increased from 100°C to 200 °C and then 300 °C, more surface oxidation is
generated, and the surface colour changes from light straw to dark blue and then powdery
grey, indicating unacceptably high levels of surface contamination. For the component
produced with forced interpass cooling, Figure 7-5(d) shows that the surface has only a small
amount of oxidation as indicated by coloration, and possibly better than that of deposition
with a naturally-cooled interpass temperature of 100 °C. The potentially better result with
forced cooling may be due to superior cooling of the zone beyond the uppermost layer. This
would allow the heat input from the newly deposited layer to be better absorbed by the
specific heat capacity of the previously deposited material, keeping the surface temperature
sufficiently low to avoid absorption of atmospheric oxygen after the shielding gas is turned
off at the end welding. These results suggest that rapid CO2 gas cooling during the dwell
interval is beneficial for avoiding post-deposit oxidation titanium when using WAAM with
localized gas shielding in an otherwise open atmosphere.
(a)
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(b)
(c)
(d)
Figure 7-5 The surface appearance of Ti6Al4V parts fabricated at different process
conditions: (a) 100°C; (b) 200°C; (c) 300°C; (d) Forced cooling with CO2 gas
7.3.3 Macrostructure
Figure 7-6 shows the cross-section macrograph of the four as-fabricated Ti6Al4V walls.
Three distinct regions are discernible on the etched cross-section: the substrate regions with
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concave bands (R3), the bottom regions with parallel bands (R1) and the top regions (R2) with
convex bands.
The concave band is apparently determined by the field of heat affected zone which
expands at high interpass temperature on the substrate. However, when the interpass
temperature is controlled at 300 °C, there are no obvious convex bands and the
microstructure produced in the substrate region is homogeneous, possibly attributing to the
large grain growth at low cooling rates. For the parallel bands and convex bands, which are
present in the majority of the deposits, their number varies with interpass temperature. Figure
7-7 further displays the proportion ratio (R1/R2) between two types of bands at different
process conditions. It can be observed that the number of parallel bands significantly decrease
with increasing interpass temperature, while the changes in convex bands is the exact
opposite. For the deposit produced using forced interpass cooling, the area of convex bands
regions is almost equal to that of parallel bands regions, as shown in Figure 7-6d and Figure
7-7. The vertical spacing of the bands in the parallel and convex regions of the deposit
produced using forced interpass cooling is visibly greater than the spacing of the bands seen
in any of the other deposits, considered to be a result of consistently higher cooling rate.
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Figure 7-6 The cross-section morphology of Ti6Al4V component fabricated at :
(a)100°C; (b) 200°C; (c) 300°C;(d) forced interpass cooling with CO2 gas
Figure 7-7 The area ratio of parallel band regions and convex band regions for different
process conditions
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7.3.4 Microstructure
Figure 7-8 shows the microstructure of the as-fabricated Ti6Al4V parts, examined along
the building direction. Due to the influences of thermal gradient and solidification rate,
near-equiaxed beta grains (Figure 7-8a) are formed in the substrate regions, lamellar α
interwoven with a coarse Widmanstätten structure are produced in the parallel band regions
(Figure 7-8b), and acicular α interwoven with a basketweave structure are generated in the
convex band regions (Figure 7-8c). Similar microstructural evolution has been reported for
fabrication using electron beam deposition [173] and laser additive manufacturing [174]. It is
well known that deposited Ti6Al4V mainly consists of α', α and β phases, and only α' phase is
formed at high cooling rates. As can be seen from Figure 7-8, the microstructures and grain
size in parallel and convex band regions are not significantly influenced by interpass
temperature, possibly due to the similar heat dissipation behaviour at the same locations
within each of the produced parts. However, for convex band regions with high interpass
temperature, the grains are obviously large and the microstructure is more uniform due to
grain growth at low cooling rates. For the part deposited using CO2 gas interpass cooling,
comparatively more acicular α phases can be observed in both convex band and parallel band
regions. That is to say, when rapid interpass cooling is applied, the large amount of lamellar α
phase would be replaced by acicular α phase, at the same time accompanied with refined
grain.
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Figure 7-8 The microstructures of parts fabricated under different process conditions
7.3.5 Hardness
Figure 7-9 shows the microhardness distribution along the vertical centerline of the
cross-section of the Ti6Al4V parts fabricated under different process conditions. There is no
obvious change in the average hardness values when the interpass temperature is increased
from 100 °C to 300 °C, but the deposition with CO2 gas cooling exhibits a slightly higher
average hardness value, which is attributed to comparatively more acicular α phase and grain
refinement. Despite the markedly different appearances of band structures in the top and
bottom regions, the microhardness is almost unchanged, suggesting that the difference in α
phase volume is not responsible for the hardness performance, as similarly explained by
Baufeld et al [67].
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(a)
(b)
Figure 7-9 Hardness profiles of specimens: (a) Hardness distribution; (b) Average hardness
7.3.6 Mechanical property
Figure 7-10 shows the ultimate tensile strength (UTS) and elongation of tensile
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specimens derived from different locations within the four parts produced under different
process conditions. The UTS is sensitive to the locations within the build, decreasing along
the build height even though the various interpass temperatures were controlled throughout
the build height, as seen in Figure 7-10a. However, the elongation of these tensile specimens
shows a slight increase and then a decrease from bottom to top of the deposit, as shown in
Figure 7-10b. From both of these results it is evident that the mechanical properties are
asymmetrical within the WAAM deposit. In addition, as a result of the combined effects from
grain size and oxidation behaviour, a significant decrease in both average tensile strength and
elongation is observed with increasing interpass temperature (Figure 7-10c). The deposit
produced with additional CO2 gas interpass cooling shows slightly higher UTS and lower
elongation compared to those deposits produced without forced cooling, suggesting that
interpass cooling is beneficial to the material properties of WAAM-fabricated Ti6Al4V
components.
(a)
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(b)
(c)
Figure 7-10 The mechanical properties of obtained specimens
7.3.7 Fracture behaviour
Figure 7-11 displays the fractographs the tested tensile samples. All the tensile samples
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present ductile morphology with a large number of shallow and stretched dimples in response
to fracture. There is no significant difference in the fracture features among all the
components that have been produced with different process conditions. However, the fracture
surface of the specimen fabricated with an interpass temperature of 300 °C presents small
amounts of lamellar structures, indicated within the red rectangle in Fig.11c, which
demonstrates rapid crack growth along the crystallographic planes.
(a) (b)
(c) (d)
Figure 7-11 High-magnification fractographs of tensile samples: (a) 100°C; (b) 200°C; (c)
300°C; (d) Interpass CO2 gas cooling
7.4 Discussion
Based on the results presented in the previous section, the influence of in-process
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interpass temperature on the deposition geometry, surface oxidation and mechanical
properties is significant, but there are no obvious changes in the microstructural evolution and
hardness values. The use of forced interpass cooling discernibly improves the material
properties, which suggests that the overall material properties of Ti6Al4V components, in this
case, may be better controlled during the WAAM process.
It is accepted practice to visually inspect the surface additively manufactured Ti6Al4V
parts for coloration and surface imperfections, as this gives an indication of the effectiveness
of the inert gas shielding envelope against atmospheric contamination [131]. In this study, the
localized gas shielding is only effective for the most recently deposited layers. Generally, the
unshielded part is easily oxidized when depositing. As found from experimental results, when
a low interpass temperature of 100°C is used, significantly less oxidation is generated on the
exposed deposition surface. The low initial component temperature results in a sufficiently
low final surface temperature (after deposition of the next layer) so the rate of oxygen
diffusion into the surface is slow enough that significant coloration does not occur after the
shielding gas flow is stopped. When the interpass temperature is increased, oxygen diffusion
into the unshielded surfaces of the part progressively increases and is visibly indicated by
surface coloration, accompanied by deteriorating bulk material properties. When using CO2
gas interpass cooling, clean fabricated surfaces are achieved, with only very light oxidation,
as indicated by a faint straw color. This demonstrates that interpass cooling is an effective
method for minimizing surface oxidation in WAAM using localized gas shielding, while also
significantly reducing the interpass dwell time. WAAM fabrication of highly reactive metals
usually involves the use of a closed chamber containing an inert gas atmosphere. This
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imposes severe constraints on equipment operation and maximum component size, as well as
the inability to introduce other production aids such as CO2 gas cooling to reduce production
times. This study proposes an effective method for fabricating reactive metals such as
Ti6Al4V using localized gas shielding in an open environment, while being able to produce
components with acceptable properties. This could play an important role in reducing
processing times and cost of production and improving manufacturing efficiency.
Although high interpass temperature for fabrication means less dwell time and hence
higher effective deposition rate, the experimental results have shown that the layer geometry
is significantly different to that produced at interpass temperatures closer to room temperature.
Unless this effect is faithfully modelled by the WAAM path planning algorithm and
compensatory steps are taken, the geometrical accuracy of the produced part will be poor.
Hence, the interpass temperature should be considered in the path planning process for
depositing and needs to be controlled during the manufacturing process. This may need to
include the initial additive deposit on the substrate. Although this aspect has not been studied
here because it would be considered as pre-heating, it may be necessary to avoid the
geometrical inaccuracies that are often produced in the first few layers. Alternatively, the
WAAM deposition parameters for the first few layers will need to have an effectively higher
ratio of heat input to produce the desired geometry.
With reference to the distribution of band structures seen in the cross-sectional
macrographs of the deposited parts, Baufeld et al. [13] have claimed that material taken from
top regions of deposits exhibiting convex bands is much weaker than material taken from
bottom regions exhibiting parallel bands. This is in agreement with the tensile results
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obtained in this study. As the interpass temperature increases, more convex bands develop,
leading to an increased area with inferior mechanical properties. However, more deposition
layers can be visible since these convex bands form in association with layered bulges which
represent the liquids of subsequent WAAM steps, as indicated by Baufeld et al. [62]. The
proportion of the deposit containing parallel bands is reduced as the interpass temperature
increases, indicating a reduction of build area with comparatively superior mechanical
properties. It should be mentioned that these parallel bands perform inconsistency with
visible deposition layers, which is evident by their end in the indentation between layered
bulges [13]. For this case, forced interpass cooling is considered as an effective way to
increase the number of parallel bands, as shown in Figure 7-6 and Figure 7-7, and the
resulting effects are even better than those for the manufacturing process with low interpass
temperature achieved from natural cooling, due to the widening of band structures at higher
cooling rates.
Referring to material properties, interpass temperature only has a small effect on the
microstructures, but it has been observed that rapid interpass cooling causes more grain
refinement, producing correspondingly better performance in hardness and mechanical
properties. The hardness of WAAM-fabricated Ti6Al4V is mainly determined by solid
solution and grain boundaries or dislocation distribution due to some segregated elements
interacting between grain boundaries, and edge dislocations existing in α/β microstructures
[73, 175]. When forced interpass cooling is used, the deposit cools down faster than for the
case of natural cooling, which thus brings more grain boundaries and dislocations to produce
higher microhardness values [145]. Furthermore, the Ti6Al4V microstructure obtained at
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high cooling rates consists of large amounts of massive α phases with acicular appearance
and non-equilibrium composition, which are usually harder and have higher strength than
those produced at lower cooling rates. For the same reason, forced interpass cooling produces
deposits with higher hardness. It should be mentioned that the massive α phase has similar
crystalline structure to lamellar α phase and has similar chemical composition to β phase
[176], which contributes to the hardness being uniformly distributed in the integral part even
when the process conditions are changed.
Another benefit of forced interpass cooling is that for Ti6Al4V α+β structures, high
cooling rate and rapid solidification rate are also expected to increase the β phase, helping to
make the phase distribution more easily approach the equilibrium (around 5% β phase and 95%
α phase) [177]. Therefore the material strength, in a manner, is enhanced.
Form the Discussion part, it is apparent that what this study prefers to display is a
comprehensive analysis of WAAM printed Ti6Al4V alloy with forced interpass cooling using
compressed CO2. It is believed that the proposed innovative process is beneficial to the
fabricated part. In near future, the best practices for applying interpass cooling and localized
gas shielding will be continued.
7.5 Conclusion
In this study, Ti6Al4V alloy has been used as the build material for the GT-WAAM
process using localized gas shielding. The effects of various interpass temperatures and CO2
gas interpass cooling on the bead geometry, surface oxidation, microstructural morphology,
grain size, mechanical properties and fracture features were investigated. The findings
include:
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(1) In terms of microstructural evolution and hardness, the properties of
WAAM-fabricated Ti6Al4V components are measurably different but not greatly affected by
variation of the interpass temperature. However, it is still necessary to control interpass
temperature to a reasonable range to ensure the geometric accuracy of each deposited layer
and to avoid surface oxidation.
(2) Forced interpass cooling using compressed CO2 gas is easily implemented, and is
beneficial to additively manufactured Ti6Al4V components, contributing to an appealing
surface finish with less visible surface oxidation, refined microstructure, improved hardness
and enhanced strength.
(3) WAAM deposition with CO2 gas interpass cooling significantly promotes
manufacturing efficiency through a sharp reduction of dwell time between the deposited
layers.
Further investigations will focus on the effects of different cooling gas flow rates and
various cooling times on the deposition properties in order to validate the effectiveness of this
process. In addition, quantitative temperature measurement using IR pyrometry will be used
to explore the thermal behaviour during deposition with interpass cooling. The residual stress
and deformation of WAAM fabricated components will also be presented in near future.
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Chapter 8 DISTORTION CONTROL
8.1 Introduction
Thermal distortion often affects every part of wire arc additively manufactured
component, since the feed materials experience non-uniform thermal expansion and
contraction under alternate heating and cooling cycles during layered deposition [141, 169].
Mukherjee et al. [178] reported the thermal distortion in additively manufactured parts is a
result of the combined effects of material properties, geometric features, deposition pathways,
process parameters and preheating or cooling conditions. Denlinger et al. [43] found that
thermal distortion is able to accumulate along the build direction due to ineffective heat
dissipation and the resultant introduced preheat for the subsequent layer. It is well known that
a WAAM-deposited component exhibits different types of distortion, including longitudinal
and transverse shrinkage, bending distortion, angular distortion and rotational distortion,
resulting in imprecise geometry of the final part and also creating residual stress that may
lead to failure during service[103]. Hence, it is necessary to control the thermal distortion to
an acceptable range to ensure adequate quality of WAAM parts.
To address this issue, it is critical to predict the thermal state or the internal stress field
changes in the deposition. A study was conducted by Cao et al.[179] to examine the distortion
characteristics and residual stress distribution in WAAM-fabricated Ti6Al4V structures, and
proposed that preheating of the substrate can significantly mitigate residual stress and
distortion. Montevecchi et al.[180] introduced a number of simulation methods to predict
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distortion during deposition, and satisfactory results were achieved by experimental
validation.
To date, both thermal and mechanical methods can be used to reduce the distortion in
WAAM-fabricated parts. Thermal methods mainly consist of heat input control and
post-fabrication heat treatment. Typically, heat input control within an acceptable range only
reduces the part distortion to a limited extent, while post heat treatment is time-consuming,
costly and limited by the size of the produced component. Attempts have been made to
develop efficient and low-cost equipment to reduce distortion, for example, using interpass
cold rolling and designing suitable manufacturing fixtures [181]. However, the interpass cold
rolling technique is only feasible for deposited parts with simple shapes such as straight walls,
due to geometrical limitations of the rolling process [114]. Designing a variety of fixtures to
minimize flexibility of work pieces obviously brings more labour and higher costs for
manufacturing [182]. Recently, active interpass cooling using compressed CO2 was applied to
the WAAM process. It was found that this innovative approach provided benefits to the
WAAM-produced Ti6AL4V alloy, contributing to an appealing surface finish with refined
microstructure, enhanced tensile strength and improved microhardness [183].
In this section, an active interpass cooling process, where compressed CO2 is used to
cool the area behind the heat source, is employed to reduce flexural distortions in WAAM
Ti6Al4V wall structures. By means of experiment and simulation, the thermal histories,
geometrical features and distortional characteristics are evaluated for components produced
with and without interpass cooling. Moreover, a comparative analysis between results is
presented, aiming to discuss the feasibility of adding sequential active interpass cooling into
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the WAAM process to control part distortion.
8.2 Experimental Procedures
8.2.1 Experiment setup
The apparatus employed in this research work as shown in Figure 8-1.
Figure 8-1 Experimental setups of the GT-WAAM deposition system with forced
interpass cooling
In this study, active interpass cooling was achieved by delivering a stable CO2 stream
through a commercially-available spray nozzle located behind the trailing gas shield cover, as
displayed in Figure 8-1. The cooling temperature achieved by the compressed CO2 was
measured at approximately -78 °C. As the relatively turbulent CO2 flow may induce arc
disruption during deposition, the CO2 spray was started immediately after the deposited layer
was completed and the arc was extinguished. Subsequently, the cooling spray was passed
over the newly deposited layer with the same speed as the deposition, and it was stopped
when the spray arrived at the end of the wall. Three nozzle types were alternatively tested,
delivering CO2 gas flow rates of 320 g/min, 730 g/min and 1460 g/min respectively. During
the build, the thermal profiles were continuously monitored by a thermocouple which was
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attached to the mid-length of the substrate and offset by 20mm from the deposition path, as
shown in Figure 8-2. The recorded data was processed using LabVIEW software.
Figure 8-2 Schematic of temperature measurement location for the thermocouple
Six straight wall structures were produced with different process parameters using the
described additive manufacturing system. Various cooling gas flow rates and cooling times
were selected to explore the influence of different cooling strategies on thermal distortion.
Accordingly, specimens were divided into three groups by cooling conditions, as listed in
Table 8-1. Group I, with test specimen coded as S1, refers to the deposition without interpass
cooling and serves as a reference. Group II with specimens from S2 to S4 and Group III with
specimens S2, S5 and S6 were designed to investigate the effects of cooling gas flow rate and
cooling time, respectively. The cooling times for Group III were determined by the ratio of
travel speeds, which compares as relative time for the investigation. The detailed parameter
design for the cooling process is provided in Table 8-1, and the deposition parameters are
listed in Table 8-2. Welding grade argon (99.995% purity) was used as shielding gas for both
welding torch and trailing shield cover.
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Table 8-1 Parameter design for cooling process
Group Deposition sample CO2 flow rate (g/min) Travel speed (mm/min) Delay time (min)
I S1 0 150 N/A
II S2 320 150 1
S3 730 150 1
S4 1460 150 1
III S5 320 230 0.65
S6 320 300 0.5
Table 8-2 Welding and deposition process parameters
Processing parameters Details
Welding current 140A
Arc voltage 15.5V
Wire feed speed 1060 mm/min
Distance from electrode to workpiece 3 mm
Angle between electrode and filler wire 60°
Flow rate of argon in GTAW welding torch 15 L/min
Flow rate of trailing shield 15 L/min
Number of deposited layer 10
8.2.2 The deformation measurement
To measure the distortion of the as-deposited specimens, a 3D laser profile scanner with
a resolution of 0.02 mm was used and the recorded data was processed with commercial
MATLAB software.
8.2.3 Modeling process
8. 2.3.1 Finite element model
As shown in Figure 8-3, only half of the symmetrical substrate is modeled in order to
reduce the simulation time. In the finite element model, the minimum mesh unit has a
dimension of 1.3 mm × 1.5 mm × 0.9 mm. ANSYS SOLID5 and a 3D coupled-field element
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with 8 nodes were adopted. SOLID5 has 3D magnetic, electric, thermal, piezoelectric and
structural field capability with limited coupling between each field. In this study, thermal and
structural field coupling was utilized. The finite element model generated 5396 units and
7804 nodes during the simulation process.
Figure 8-3 3D finite element mesh for one half of build plate
8.2.3.2 Heat source
In the numerical simulation, the heat source distribution is described by the following
expression:
𝑞(𝑥, 𝑦) = 3𝑃 ∙ exp (−3(𝑥2 + 𝑦2)/𝑅2)/(𝜋𝑅2) (8-1)
where q(x, y) is heat flux density; P is the power of welding arc; η is the welding arc
efficiency (η =0.7); x2 + y2 is the square of the distance from the calculation point to the
center of the arc and R is the radius of the arc ( R =10 mm).
8.2.3.3 Cooling process
When simulating the air-cooled deposition process (Group I), the temperature at the
reference point would be checked at the end of each step. Each cooling step lasted for 100
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seconds. Once the temperature of the reference point reached room temperature, the natural
cooling process would be terminated and the welding process for the next layer would be
initiated.
In the simulation of deposition with CO2 interpass cooling (Group II and Group III),
nodes at the CO2 spray cooling position were applied with temperature constraints using the
data measured respectively after each cooling process to simulate the rapid temperature drop
of the weldment. As the gas nozzle moved along the welding direction, the cooling position
changed with time. Therefore, at the beginning of each step, temperature constraints of the
position that had just been cooled down in the previous step would be removed and new
temperature constraints would be added to the new cooling region.
8.2.3.4 Boundary conditions and material attributes
A reference temperature was set to indicate the initial temperature at the beginning of the
simulation. Reference temperature was also used to simulate room temperature during the
calculation. Displacement constraints that simulate the fixture of the workpiece have great
influence on the stress calculation and displacement results. Four corners of the plate were set
to rigid displacement in all three directions to simulate the constraints of the fixture bolts
during the deposition process, as shown in Figure 8-4. After process completion, the
constraints at these nodes were deleted to simulate the removal of the bolts after the process.
Also, in order to avoid drift of the entire structure, one corner node of the plate was set to full
constraint and three nodes at the other three corners were set to rigid displacement in the
vertical direction. Displacements parallel to the transverse direction were constrained on the
symmetrical face.
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Figure 8-4 The boundary conditions for the simulated model
Temperature field simulation of the welding process is a non-linear transient analysis.
Properties of the material, such as yield strength, elastic modulus, specific heat and thermal
conductivity are temperature-dependent. Therefore, property data is associated with a
temperature table. Poisson's ratio is considered to be constant since the change of Poisson's
ratio with temperature is quite small. Therefore, it was set to 0.3 during the whole simulation
process.
The workpiece was encompassed by varied thermal mediums. Heat transfer from the top
surface of the workpiece was mainly conducted by air, and heat transfer from the bottom
surface was considered to be forced convection to the worktable. As a result, heat transfer
coefficients were set differently in different areas. No heat transfer coefficient was set on the
symmetrical face.
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8.3 Results and discussion
8.3.1 Thermal behaviours during deposition
Figure 8-5 shows the thermal profiles measured by thermocouple during the deposition
process, with and without active interpass cooling. It is observed that when the welding torch
moved along the travel direction, the temperature increased rapidly until a peak temperature
was achieved. After completing the deposition of each layer, the temperature gradually
decreased with natural cooling, but a much sharper drop in temperature occurs with interpass
cooling. As a result, the build time for the 10 deposited layers in this study was reduced by at
least 80 percent when CO2 cooling was implemented, as comparably shown in Figure 8-5. It
is evident that active interpass cooling in WAAM can offer a significant build-time saving
through a dramatic reduction of interlayer dwell time, and thereby provides a means for
improvement of manufacturing efficiency.
(a)
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(b)
Figure 8-5 The temperature profiles of deposition (a) with natural cooling and (b) with active
interpass cooling
Variations in thermal profile occur when the CO2 cooling gas flow rate or cooling time
are changed during fabrication. Figure 8-6 compares the thermal histories of one selected
layer in all five depositions using different cooling parameters. With increased CO2 gas flow
rate from 320g/min (S2) to 1460g/min (S4), the peak temperature could be reduced by over
50% due to a significant increase in heat removal from the component after the previous layer
was deposited. As well, this trend can be observed in processes with increased travel speeds
(S2, S5 & S6), whereby lower heat input brings likewise lower heat accumulation along the
build direction. Hence, it is accepted that the combined effects of heat input, cooling gas flow
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rate and cooling time potentially determines the development of peak temperature, which
could be related to the different cooling effects in the deposition. Additionally, active
interpass cooling can cause uneven distribution of heat in the build part. Due to the increased
thermal dissipation in the uppermost layer of deposit, the interior part produces a large
difference between the highest and the lowest values that cooling rates may take. As the
deposition suffered from uneven heat distribution, it expands irregularly in all directions,
resulting in inhomogeneous stress and distortion to the build part.
Figure 8-6 The thermal history of one selected layer with different active interpass
cooling processes for specimens S2 to S6
8.3.2 Geometrical features
Figure 8-7 shows the geometrical features of as-deposited specimens in this study.
Compared to the specimen produced with natural cooling (S1), when using active interpass
cooling at the same travel speed (S2, S3 and S4) the average width of build walls was
reduced by around 7 % and average height increased by about 6 %. Due to the rapid cooling
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produced by compressed CO2, a decreasing amount of heat distributed in the previous
deposition is considered as the part of heat input for the subsequent layer, contributing to
lower temperatures in the molten pool and surrounding area, and thus a narrower layer weld
bead. So, although active interpass cooling can be a tremendous advantage to promote
manufacturing efficiency, the results have shown that the build geometry changes slightly.
Unless, this effect is correctly modelled in path planning step or compensatory steps are taken
during the build process, the geometrical accuracy of the final part will be compromised.
Therefore, when the proposed cooling method is implemented in WAAM, it should also be
considered in the path planning strategy to ensure geometrical accuracy of the product.
It should be noted that specimens S5 and S6 were produced at different deposition travel
speeds, so the bead dimensions would be different to S1, even before the additional influence
of active interpass cooling.
Figure 8-7 Effect of forced interpass cooling on wall dimensions
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8.3.3 Thermal distortion
Figure 8-8 shows the simulation results for the distortion of the additively manufactured
specimens. Figure 8-9 further displays the distortion of the substrate, examined along the
longitudinal direction (travel direction) and transverse direction (perpendicular to the travel
direction). When adding sequential interpass cooling to the deposition process, the produced
part sees an 81% and 69% in the maximum reduction of the longitudinal and transverse
distortion, respectively. The simulation results are basically consistent with experimental
observation, which indicates that active interpass cooling is able to significantly mitigate the
distortion issues in WAAM parts.
(a) S1 (b) S2
(c) S3 (d) S4
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(e) S5 (f) S6
Figure 8-8 Simulation results for the distortion of produced samples: (a) S1; (b) S2; (c) S3;
(d) S4; (e) S5; (f) S6.
(a) Longitudinal distortion
(b) Transverse distortion
Figure 8-9 The thermal distortion of as-deposited specimens: (a) longitudinal distortion and
(b) transverse distortion of half of substrate.
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Due to non-uniform thermal contraction in the melt area, shrinkage, which is inhibited by
underlying materials, takes place. The generation and development of shrinkage induces a
tensile stress in the newly added top layer and a compressive stress is generated in underlying
layers, which leads to a localized distortion in the deposited areas. During deposition, theses
localized distortions, particularly in the first few layers, accumulate along the build direction.
These lead to distortion and dimensional variation in the final part, as shown in Figure 8-10.
In extreme cases, complete loss of the planned bead deposit dimensional control may occur. It
is worth noting that part distortion reached a steady-state condition after several layers of
build, due to the thermal strain reaching a nearly constant value. Obvious distortion is
generated along the travel direction, attributable to the large thermal strain.
Figure 8-10 The distortion mechanism in WAAM-fabricated Ti6Al4V process
Two main parameters, cooling gas flow rates and cooling time, are primarily responsible
for the change in distortion of parts produced using the WAAM process with active interpass
cooling. Mukherjee et al. reported that that thermal distortion of an additively manufactured
part is dependent on a strain parameter, ε*[184]:
CHAPTER 8 DISTORTION CONTROL
139
ε ∗=𝛽∆𝑇𝑡𝐻
32
𝐸𝐼𝐹√𝜌 (8-2)
𝐹 =𝛼𝜏
𝑤2 =𝛼
𝑉𝐿 (8-3)
where β is the volumetric coefficient of thermal expansion, E is the elastic modulus, ΔT is the
maximum rise in temperature, I is the moment of inertia of the substrate, t is the characteristic
time, H is the heat input per unit length, F is the Fourier number which is the ratio of the heat
dissipation rate to heat storage rate, ρ is the density of the alloy wire, α refers to thermal
diffusivity, τ is characteristic time scale, w is deposition length, V is the traveling speed, and L
is the length of molten pool. Eq (2) and (3) demonstrate that both reducing heat accumulation
and improving heat dissipation in a deposited part can result in a high Fourier number and
low peak temperature, leading to a low thermal strain and distortion for the additively
manufactured part [178]. With the assistance of active interpass cooling, the heat stored in the
deposit can be quickly dissipated, contributing to less heat accumulation. This is useful for
the reduction of distortions that are difficult to mitigate with mechanical methods. In addition,
distortion accumulation can also be reduced by decreasing the interlayer dwell time as
pointed out by Denlinger et al. When active interpass cooling used, the inherently shortened
dwell period brings less distortion accumulation, contributing to the desired geometrical
performance of the final produced part.
In comparison to the results produced by solely adjusting the cooling gas flow rate
(Group II), the effect of adjusting both the deposition speed and the cooling time (Group III)
provides improved mitigation of distortion, which may be attributed to the combined effects
of reduced heat input at deposition and even further shortened dwell time between deposits,
CHAPTER 8 DISTORTION CONTROL
140
creating greater average heat dissipation per unit time during the build process. From this
point of view, if active interpass cooling is used, distortion control is more effective through
setting the cooling time rather than purely varying the cooling gas flow rate. By this approach,
both cooling efficiency improvement and cooling gas savings can be achieved. However, the
shape of the deposited material is necessarily different, so the path planning algorithm must
compensate for such changes.
The primary focus of this paper has been to propose and evaluate an innovative method
for low distortion WAAM fabrication of metal components, while promoting high
manufacturing efficiency. Active interpass cooling using compressed CO2 is effective in
reducing distortion of WAAM components and improving dimensional accuracy. Future work
will investigate the residual stress distribution in components which are produced by different
cooling times and cooling gas flow rates during fabrication.
8.4 Conclusion
In this study, an innovative WAAM process with active interpass cooling using
compressed CO2 was developed for fabricating Ti6Al4V wall structures. The thermal
behavior, geometrical features, and distortion characteristics of as-fabricated specimens were
investigated and discussed. The following conclusions could be drawn:
(1) By applying active interpass cooling between deposition of layers, the build times
have been reduced by at least 80%.
(2) By means of active interpass cooling, significant distortion reduction is possible in
produced parts due to rapid heat removal and consequently decreased heat accumulation
CHAPTER 8 DISTORTION CONTROL
141
during deposition. It was observed that simultaneously altering deposition and cooling times
has a greater mitigation effect on thermal distortion than only adjusting the cooling gas flow
rate, which is possibly related to the difference in the overall heat dissipation per unit time.
(3) Since active interpass cooling affects the deposition process, the path planning
strategy should take this into consideration, and more optimized strategies for distortion
control should be considered from the cooling time point of view.
CHAPTER 9 CONCLUSION
142
Chapter 9 CONCLUDING REMARKS
Through matching a knowledge of material characteristics with the performance features
of particular WAAM techniques, a quality-based framework is proposed, for producing
high-quality and defect-free components in this chapter. A brief summary and future research
direction are following provided.
9.1 A quality-based framework for WAAM process
Improving process stability, eliminating or decreasing deposition defects and producing
components with high quality and mechanical performance have become major research
focuses in making the WAAM process more competitive against other additive
manufacturing methods. An in-depth understanding of various materials, ideal process setup,
in-process parameter control and post processing is essential for achieving such a goal. After
a systematic review and analysis, a quality-based framework aiming to achieve high-quality
and defects-free WAAM process is proposed, as shown in Figure 9-1. Three main aspects are
considered: feedstock optimization, manufacturing process, and postprocess treatment.
CHAPTER 9 CONCLUSION
143
Quality performance
Process selection Feedstock optimization Post-process treatment
TIG
• Parameters coordination
• Gas mix
MIG
• Process type ( CMT AC etc)• Gas mix
Plasma
• Heat input• Gas mix• Hot wire
Process monitoring & control
• Deposition geometry• Interpass temperature• Metal transfer behaviour
Defects control
• Geometrical accuracy
• Porosity
• Cracking, delamination & swelling
• Oxidation
Wire design
• Standard wire• Wire mixing• Development wires
Heat treatments
• Ageing• Solution treatments• Combinations
Work hardening
• Peening• Ultrasonic • Cold-rolling
Microstructural evolution
In-process treatment
• Replanning path • Interpass cold rolling• Interpass gas cooling• Ultrasonic • Surface finishing
Material properties
• Strength & ductility• Fatigue & damage
tolerance• Corrosion property
Quality improvement
• Distortions• Residual stress• Microstructure• Hardness
Interpass gas cooling
• Cooling gas• Cooling time• Cooling methods
Figure 9-1 A quality-based framework for WAAM process
Selection of the most suitable welding WAAM process for the deposition material can
ensure manufacturing process stability and contribute to reduction of defects. For example, if
the CMT-PADV process is used for producing aluminum parts, porosity defects can be
dramatically reduced when compared to other GMAW modes [98]. Moreover, integrated and
reliable process monitoring and control systems are needed to maintain the stability of the
process and ensure the quality of production. Usually, the bead geometry, interpass
temperature, arc characteristics and metal transfer behaviour are included in process
monitoring and control. Controlling the interpass temperature within a reasonable range is
beneficial to microstructural evolution and the resulting mechanical properties. Further,
regulating the arc characteristics and metal transfer behaviour in real time is helpful to
process stability and avoidance of defects. Based on the process monitoring data that has
been collected during deposition, one (or more) of several postprocess treatments can be
CHAPTER 9 CONCLUSION
144
selected to mitigate defects and improve mechanical performance.
Considering the material characteristics, microstructural evolution and mechanical
properties can also be optimized through new feedstock composition design. It is well known
that different alloying elements have specific effects on material properties. By referring to
the phase diagram, the desired deposition microstructures can be obtained via adding specific
alloying elements in the feedstock and subsequent mechanical properties improved. Moreover,
availability of wire mixings provides a potential possibility for building large functionally
graded products for special applications. For example, twin-wire GTAW-based WAAM has
been successfully developed to produce intermetallic graded materials [24, 90]. The
development of new powder cored wires will offer exciting opportunities for fabricating
target components with accurate metal composition. In summary, using new welding
consumables brings a cost-effective solution, which supports high deposition quality through
obtaining the desired microstructures, lowering manufacturing costs by reducing or
eliminating pre-weld cleaning and re-work, and providing safer working environments by
reducing weld fumes.
Another essential part of WAAM processing for most materials is post-process treatment,
which is used to reduce residual stresses and distortion, refine microstructures, improve
microhardness and enhance material strength. However, post-process technologies have their
own limitations, for instance, peening and ultrasonic impact treatment only improve material
property and reduce defects near the part surface, while extended heat treatment of certain
materials promotes grain growth rather than grain refinement. Currently, most WAAM
fabricated parts need to be post-process treated with a selective combination of technologies
CHAPTER 9 CONCLUSION
145
to reduce the defects and improve the product quality to the greatest extent possible.
9.2 Future perspective
In WAAM of metallic materials, the fundamental interrelationships between material
composition and microstructure govern the material properties and fabrication quality. Since
the WAAM process is an inherently non-equilibrium thermal process, it is challenging to
predicate and control the microstructural evolution, which is responsible for the variation of
mechanical properties in the deposited part. Further research attention should be paid on the
study of underlying physical and chemical metallurgical mechanisms in WAAM process to
provide a guidance for the process optimization, improvement and control. The defects
generated in WAAM-produced part are closely related to the target material characteristics
and process parameters. The development of strategies or ancillary process to overcome
defects generation are of prime importance. With the requirement of high quality WAAM part,
the proposed quality-based framework will see a wide application in the future years.
As WAAM matures as a commercial manufacturing process, development of a
commercially available WAAM system for metal components is an interdisciplinary
challenge, which integrates physical welding process development, materials science and
thermo-mechanical engineering, and mechatronic and control system design. Although much
research has been carried out over recent years in various areas, such as process planning,
programming, and material study, a general purpose off-the-shelf WAAM system similar to
commercially available powder-bed fusion systems is yet to be developed. Due to the variety
of requirements of different engineering materials and the varying scale of fabrication, many
different WAAM system designs are expected to be developed that will be optimized for
CHAPTER 9 CONCLUSION
146
particular applications, rather than a single system that is capable of addressing all of the
possible problems.
PUBLICATION
147
PUBLICATION
[1] Bintao Wu, Zengxi Pan *, Donghong Ding, Dominic Cuiuri, Huijun Li, Jing Xu, John
Norrish. A review on the wire arc additive manufacturing of metals: property, defect and
quality improvement. Journal of Manufacturing Processes, 35 (2018):127-139. – from
Chapter 2
[2] Bintao Wu, Donghong Ding, Zengxi Pan*, DominicCuiuri, Huijun Li, Jian Han, Zhenyu
Fei. Effects of heat accumulation on the arc characteristics and metal transfer behavior in
Wire Arc Additive Manufacturing of Ti6Al4V. Journal of Materials Processing
Technology, 250 (2017):304 -312. – from Chapter 4
[3] Bintao Wu, Zengxi Pan*, Donghong Ding, Dominic Cuiuri, Huijun Li. Effects of heat
accumulation on microstructure and mechanical properties of Ti6Al4V alloy deposited
by wire arc additive manufacturing. Additive Manufacturing, 23(2018):151-160. – from
Chapter 5
[4] Bintao Wu, Zenxi Pan*, Siyuan Li, Huijun Li,Donghong Ding, Dominic Cuiuri. The
anisotropic corrosion behaviours of wire arc additive manufactured Ti-6Al-4V parts in
3.5% NaCl solution. Corrosion Science, 137 (2018):176 -183. -from Chapter 6
[5] Bintao Wu, Zengxi Pan *, Donghong Ding, Dominic Cuiuri, Huijun Li, Zhenyu Fei. The
effects of forced interpass cooling on the material properties of wire arc additively
PUBLICATION
148
manufactured Ti6Al4V alloy, Journal of Materials Processing Technology, 258
(2018):97-105. -from Chapter 7
[6] Bintao Wu, Zengxi Pan *, Donghong Ding, Guangyu Chen, Lei Yuan, Huijun Li,
Dominic Cuiuri, Stephen van Duin. Mitigation of thermal distortion in wire arc additively
manufactured Ti6Al4V parts using active interpass cooling. Science and Technology of
Welding and Joining, 2018, (Accept)-from chapter 8
[7] Bintao Wu, Zengxi Pan *, Donghong Ding, Huijun Li. Thermal behaviour in Wire arc
additive manufacturing: characteristics, effects and control. Transactions on Intelligent
Welding Manufacturing, 2018, (Accept).
[8] Bintao Wu, Zengxi Pan *, Huijun Li, Donghong Ding, Dominic Cuiuri, Lei Yuan, Stephen
van Duin. Innovative fabrication of aluminum pillar structure using droplet based wire
arc additive manufacturing. Journal of manufacturing process and technology, (Under
review).
[9] Bintao Wu, Zengxi Pan *, Huijun Li, Donghong Ding, Dominic Cuiuri. The effects of gas
cooling flow rate and time on the microstructure and mechanical properties of wire arc
additively manufactured Ti6Al4V alloy. Material Science and Engineering A. (Prepare to
submit).
[10] Bintao Wu, Zengxi Pan *, Huijun Li, Donghong Ding, Dominic Cuiuri. Investigation on
the microstructures and mechanical properties of wire arc additively manufactured Fe/Ni
functional part. Material & Design (Prepare to submit).
PUBLICATION
149
[11] Zengxi Pan, Donghong Ding *, Bintao Wu, Dominic Cuiuri, Huijun Li, John Norrish.
Arc Welding Processes for Additive Manufacturing: A Review. Transactions on
Intelligent Welding Manufacturing, 2017, pp:3-24.
[12] Jian Han, Cheng Lu, Bintao Wu *, Jintao Li, Huijun Li, Yao Lu, Qiuzhi Gao. Innovative
analysis of Luders band behaviour in X80 pipeline steel. Materials Science and
Engineering A, 683 (2017):123–128.
[13] Yugang, Miao, Benshun, Zhang, Bintao Wu*, Xiaoxiao Wang, Guangyu Chen,
Duanfeng Han. Joint characteristics and corrosion properties of bypass-current
double-sided Arc-welded aluminum 6061 alloy with Al-Si filler metal. Acta Metallurgica
Sinica (English Letters), 29 (2016):360 –366.
[14] Zhengyu Fei, Zengxi Pan, Dominic Cuiuri, Huijun Li, Bintao Wu, Donghong Ding,
Lihong Su, AA. Gazder. Investigation into the viability of K-TIG for joining armour
grade quenched and tempered steel. Journal of Manufacturing Processes, 32
(2018):482-493.
[15] Zhengyu Fei, Zengxi Pan, Dominic Cuiuri, Huijun Li, Bintao Wu, Lihong Su,
Improving the weld microstructure and material properties of K-TIG welded armour steel
joint using filler material. The International Journal of Advanced Manufacturing
Technology, 10 (2018): 1-14.
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